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STRUCTURE/PROPERTY RELATIONS OF 
ELASTOMERIC HYBRID ORGANIC-INORGANIC COMPOSITES 



By 
THOMAS M. MILLER 



A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL 

OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT 

OF THE REQUIREMENTS FOR THE DEGREE OF 

DOCTOR OF PHILOSOPHY 



UNIVERSITY OF FLORIDA 

1997 









To Yvonne 






ACKNOWLEDGMENTS 

Sincere thanks go to my advisor Dr. Anthony Brennan for providing me the 
opportunity to educate myself in the field of polymer engineering. The challenges 
associated with working in his group and the enthusiasm he possesses for both his students 
and his profession have made my studies at UF enjoyable and rewarding. I would also like 
to thank Dr. Christopher Batich, Dr. Elliot Douglas, Dr. Eugene Goldberg and Dr. 
Kenneth Wagener for serving on my doctorate committee. Partial financial support for 
this research was made possible through a DuPont Young Faculty Investigator Award 
presented to Dr. Anthony Brennan. Additionally, support was provided by the National 
Institutes of Health under Grant DEO 9307-07, as well as a research grant awarded by 
PhotoSense LLC. As I have benefited from all of these contributions I gratefully 
acknowledge the support of DuPont, NIH and Dr. Alan Baron of PhotoSense. To this 
last individual I am especially grateful for both his enthusiasm as a project sponsor and the 
technical insights he has provided regarding the work detailed in Chapter 5. Also 
appreciated is the financial support of the Department of Materials Science and 
Engineering provided me during my first year of study. This assistance and the air of 
professionalism present within the MSE department were instrumental in my decision to 
attend the University of Florida. 

Professional thanks go out to the members of Dr. Brennan's research group. I am 
indebted to Rick Feller and Yiqun Wang for their patience in explaining the operation of 



in 



laboratory equipment when I arrived, and for encouraging my own personal development 
with the occasional lack of help. Professional thanks are extended to Jesse Arnold and 
Mike Zamora for their unselfish donation of AFM skill and chemistry expertise, 
respectively. Additionally, their participation in numerous discussions has contributed 
immeasurably to the success of my research efforts. Also appreciated is the assistance of 
Licheng Zhao in the synthesis of the lumiphore containing samples evaluated in Chapter 5. 
The collection of the small angle X-ray scattering data presented in Chapter 4 was 
obtained by Dr. Jianye Wen at Virginia Polytechnic Institute and State University. His 
assistance is appreciated. Professional thanks are also extended to Dr. Bruce Carroll, Dr 
Paul Hubner and Andy Winslow of the Department of Aerospace Engineering, Mechanics 
and Engineering Science at the University of Florida for their assistance in collecting the 
data necessary to measure the diffusivity of the hybrids. This work is detailed in Chapter 5. 
Personal thanks go out to all unmentioned members of the Brennan group and the 
remaining MAE students, for helping make graduate school fun. I thank my parents, 
Wayne and Peggy Miller, for their lifelong commitment to my education. Lastly, I thank 
my wife, Yvonne, for her patience as I completed these last few years of study. 









IV 



TABLE OF CONTENTS 

page 

ACKNOWLEDGMENTS iii 

LIST OF TABLES viii 

LIST OF FIGURES ix 

ABSTRACT xiii 

CHAPTERS 

1 INTRODUCTION 1 

2 THE DEVELOPMENT AND APPLICATION OF HYBRID COMPOSITES 6 

Development and Detailing of Specific Composite Systems 6 

Poly(Dimethyl Siloxane) 7 

Poly(Tetramethylene Oxide) 10 

Other Rubber Based Systems 11 

Polyoxazolines 12 

Vinylic Based Polymers 13 

High Performance Thermoplastics 17 

Nafion® Membranes 18 

Clay Based Hybrids 19 

Applications Utilizing Hybrid Technology 21 

Microelectronics Usage 22 

Abrasion Resistant Coatings 22 

Solid Polymer Electrolytes 23 

Crosslinking Agents 23 

Nonlinear Optical Materials 24 

Luminescent Based Instrumentation and Devices 26 

Air/Gas and Liquid Separation Membranes 29 

Diffusion in porous materials 30 

Diffusion in polymeric materials 37 

Hybrid organic-inorganic composite membranes 41 

3 RUBBER ELASTICITY EVALUATED USING DYNAMIC MECHANICAL 
SPECTROSCOPY AND EQUILIBRIUM SWELLING 44 



Relevant Background 44 

Experimental .48 

Results and Discussion 52 

Physical Characteristics 52 

ATR-FTER Spectroscopy 53 

DMS and Estimation of Average Molar Mass Between Crosslinks 60 

Equilibrium Swelling and Estimation of Average Molar Mass Between 

Crosslinks 66 

Conclusions 68 

4 STRUCTURE/PROPERTY BEHAVIOR OF ORGANIC-INORGANIC 
SEMI-IPNs: EFFECT OF POLYSILICATE LOADING AND CO- 
SOLVENT SYSTEM 71 

Relevant Background 71 

Experimental 75 

Results and Discussion 79 

Physical Characteristics 79 

FTIR Spectroscopy 80 

Effect of Radiation on Mechanical Tensile Response 90 

Monomer Swelling and Polymer Formation 93 

Mechanical Response of PMAA-PTMO SIPNs 96 

Dynamic Mechanical Response of the SIPNs 102 

Preliminary Investigations into SIPN Morphology 106 

Conclusions 110 

5 THE EFFECT OF SOL CATALYST UPON OXYGEN DIFFUSION 1 12 

Relevant Background 112 

Motivation 112 

Enabling Principle 113 

Mass Transport Equation Utilized 115 

Catalyst Effects on Polysilicate Porosity 117 

Experimental 119 

Results and Discussion 126 

Film Thickness Measurements 126 

Dynamic Mechanical Spectroscopy 127 

Oxygen Diffusivity 129 

Conclusions 145 

6 CLOSING REMARKS 147 

Rubber Elasticity and Nonideal Networks 147 

High Performance SIPNs 149 

Gas Transport in Hybrid Composites 151 



VI 



LIST OF REFERENCES 154 

BIOGRAPHICAL SKETCH 163 



VII 






LIST OF TABLES 
Table page 

3 . 1 Room temperature densities of the TEOS(40) gels investigated 52 

3.2 Number of elastically active network chains per unit volume at each frequency 

measured using the dynamic mechanical spectrometer. N v x 10" 27 

(chains/m 3 ) 63 

4. 1 Densities of the benchmark and y-irradiated gels, as well as the estimated 

volume of polysilicate based upon the benchmark PTMO-polysilicate gel 
densities and calculations similar to those of Huang et al 79 

4.2 Primary absorbances occurring in the infrared for gel derived silica 84 

4.3 The effect of polysilicate loading upon the stress and elongation at failure of the 

PMAA-PTMO SIPNs for both co-solvent systems 99 

4.4 The effect of polysilicate loading upon the tensile yield stress and elongation and 

post-yield (P.Y.) response of the PMAA-PTMO SIPNs for both co-solvent 
systems 100 

5.1 Estimated thickness for the samples used in oxygen diffusivity measurements 126 

5.2 Diffusion model parameters and resulting oxygen diffusivity 138 

5.3 Tortuosity factors as a function of increasing polysilicate loading for the HC1 

catalyzed gels 144 



Vlll 



LIST OF FIGURES 
Figure gage 

2. 1 Schematic illustration of the relative change in diffusive flux accompanying a 

change in pore diameter for a gas of constant mean free path 31 

2.2 The pressure dependence of the mean free path of oxygen at 0°C and 100°C, 

the temperature extremes most likely encountered for any oxygen sensing 
application 34 

3.1 Reaction schematic illustrating the synthesis of end fimctionalized poly(tetra- 

methylene oxide) via the reaction of a 2% molar excess of 

isocyanatopropyl-triethoxy-silane and 2,000 g/mole poly(tetramethylene 

ether) glycol 49 

3.2 Low wavenumber region of the ATR-FTIR spectra of a benchmark TEOS(0) 

gel, benchmark TEOS(40) gel and TEOS(40) gels exposed to the 

ethylamine water solution for the indicated times 54 

3.3 High wavenumber region of the ATR-FTIR spectra of a benchmark TEOS(0) 

gel, benchmark TEOS(40) gel and TEOS(40) gels exposed to the 

ethylamine water solution for the indicated times 55 

3.4 The ratio of peak intensities for the asymmetric Si-O-Si and Si-OH stretches 

(1050/955 cm" 1 ) as well as the ether linkage of the PTMO and Si-OH 

stretch (1100/955 cm" 1 ) 57 

3.5 Low wavenumber region of the ATR-FTIR spectra of the TEOS(40) gels 

exposed to the ethylamine water solution for the indicated times after 

subtracting out the benchmark TEOS(0) spectra. The spectra, therefore, 

are those of the polysilicate phases in the TEOS(40) hybrids 58 

3.6 Dynamic mechanical storage modulus, E', as a function of temperature for the 

indicated gels. Note the different storage modulus range for the TEOS(0) 

gel, which was done to clearly show the crystallization observed 61 

3.7 The rubbery regime of the gels investigated expressed in terms of thermal 

energy. The slope of each line is the number of elastically active network 
chains per unit volume, N v 62 



IX 



3.8 Frequency dependence of the average molar mass between crosslinks, M c , for 

the indicated gels 64 

3.9 Values of the equilibrium volume fraction of polymer present in the swollen 

network, V2m, for TEOS(40) gels as a function of ethylamine solution 
exposure time 67 

3.10 Comparison of the values for M c calculated using Equations 3.5, 3.6 and 3.7 

based on equilibrium swelling using two values of Xu> as we ^ as values for 

M c obtained using Equations 3.4 and 3.5 for the 0. 1 Hz and 10 Hz data 

from the DMS 69 

4. 1 Transmission FTIR spectra for poly(tetramethylene ether) glycol and the 

subsequent triethoxysilane functionalized poly(tetramethylene oxide) 81 

4.2 ATR-FTIR spectra of the silicate fingerprint region of the THF-EPA and DMF- 

IPA benchmark gels at the indicted polysilicate loadings 83 

4.3 The ratio of the asymmetric Si-O-Si stretch (1040 cm" 1 ), Si-OH stretch (960 cm' 

) and Si-OH rocking/siloxane ring stretching (565 cm" 1 ) to the ether stretch 
attributed to the PTMO (1 100 cm' 1 ) as a function of the polysilicate loading 
for both co-solvent classes 85 

4.4 ATR-FTIR spectra of the polysilicate phases present in the gels synthesized 

from sols employing tetraethoxysilane. Spectra were obtained by 

subtracting the 4.5 vol.% gel spectra from each of the three higher 

loadings 87 

4.5 The ratio of Si-OH rocking/siloxane ring stretch (565 cm-1), Si-OH stretch 

(960 cm" ) and high wavenumber shoulder of the primary absorbance band 

in silica (1 196 cm" 1 ) to the characteristic asymmetric Si-O-Si stretch at 

1040 cm" 1 as a function of polysilicate loading for both classes of gels 89 

4.6 Effect of polysilicate loading upon the tensile mechanical response of the 

benchmark and y irradiated PTMO-silica hybrids exposed to 350 rads/min 

for a total dose of 0.069 Mrads. The age of the THF-IPA gels was 29 

days, while the age of the DMF-IPA gels was 52 days 92 

4.7 The effect of polysilicate loading upon the equilibrium MAA absorption and 

PMAA SIPN formation for the benchmark PTMO-silica hybrids 94 

4.8 The effect of polysilicate loading upon the tensile mechanical response of 29 day 

old THF-based PMAA-PTMO SIPNs 97 

4.9 The effect of polysilicate loading upon the tensile mechanical response of 52 day 

old DMF-based PMAA-PTMO SIPNs 98 



4. 10 Post-yield stress drop and elongation prior to strain hardening as influenced by 

polysilicate loading and co-solvent system employed 101 

4. 1 1 Dynamic mechanical storage modulus, E', for both the THF- and DMF-based 

PMAA-PTMO SIPNs of the same age as a function of temperature for the 
polysilicate loadings indicated 103 

4. 12 Dynamic mechanical tan 6 response for both the THF- and DMF-based 

PMAA-PTMO SIPNs of the same age as a function of temperature for the 
polysilicate loadings indicated 104 

4.13 Atomic force microscopy image of the fracture surface of a 19 vol % percent 

polysilicate, THF-based benchmark PTMO-silica gel collected using 

tapping mode 107 

4. 14 Atomic force microscopy image of the fracture surface of an 12 vol.% percent 

polysilicate, THF-based, PMAA-PTMO SIPN gel collected using tapping 
mode. This SIPN is formed from the same "parent" PTMO-silica gel as the 
piece imaged in Figure 4. 13 108 

4. 15 Small angle X-ray scattering profiles of the 19 vol.% polysilicate, THF-based 

benchmark gel and the subsequent 12 vol % polysilicate, THF-based 
PMAA-PTMO SEPN resulting from y polymerization of the MAA swollen 
"parent" gel 109 

5. 1 Schematic illustration of the types of porosity anticipated to be present as a 

result of the catalysts employed in this study 118 

5.2 Schematic illustration of the apparatus used to measure the time dependent 

intensity of luminescent gel samples 123 

5.3 Dynamic mechanical storage modulus and tan 6 response of gels containing 33 

vol.% polysilicate derived from sols employing acid and base catalysts 128 

5.4 Step responses of the four HC1 catalyzed gels utilizing the in-situ precipitation 

of polysilicate in the presence of 100% of the stoichiometric water required 
for hydrolysis 130 

5.5 Step responses of the four HC1 catalyzed gels utilizing prehydrolysis of TEOS in 

the presence of 100% of the stoichiometric water required for hydrolysis to 
produce polysilicate clusters prior to sol batching 131 

5.6 Step responses of the four ethylamine catalyzed gels utilizing in-situ 

precipitation of polysilicate in the presence of 47% of the stoichiometric 

water required for complete hydrolysis 132 



XI 



5.7 Step responses and diffusion model curve fits for HCL and ethylamine catalyzed 

gels containing PtTFPP and an estimated 4.5 vol.% polysilicate 134 

5.8 Step responses and diffusion model curve fits for the in-situ precipitated and 

prehydrolyzed HCL catalyzed gels, as well as the ethylamine catalyzed gels. 
All samples contain PtTFPP and an estimated 1 1 vol.% polysilicate 135 

5.9 Step responses and diffusion model curve fits for the in-situ precipitated and 

prehydrolyzed HCL catalyzed gels, as well as the ethylamine catalyzed gels. 
All samples contain PtTFPP and an estimated 19 vol.% polysilicate 136 

5.10 Step responses and diffusion model curve fits for the in-situ precipitated and 

rehydrolyzed HCL catalyzed gels, as well as the ethylamine catalyzed gels. 

All samples contain PtTFPP and an estimated 42 vol.% polysilicate 137 

5.11 Oxygen diffusivity as a function of polysilicate loading for all three classes of 

gels produced 140 



xn 



Abstract of Dissertation Presented to the Graduate School 
of the University of Florida in Partial Fulfillment of the 
Requirements for the Degree of Doctor of Philosophy 

STRUCTURE/PROPERTY RELATIONS OF 
ELASTOMERIC HYBRID ORGANIC-INORGANIC COMPOSITES 

By 

Thomas M. Miller 

December 1997 

Chairman. Anthony B. Brennan 

Major Department: Materials Science and Engineering 

Hybrid organic-inorganic composites have been synthesized by the sol-gel 

processing of triethoxysilane end functionalized poly(tetramethylene oxide) and 

tetraethxoysilane. The resulting transparent materials are elastomeric gels crosslinked by 

an amorphous polysilicate phase. Elementary rubber-elasticity theory in conjunction with 

dynamic mechanical spectroscopy was applied to these seemingly nonideal networks to 

quantify the change in phase interaction induced by aging the benchmark acid catalyzed 

gels in a basic solution of 70% ethylamine in water. The change in the average molar mass 

between crosslinks explained the previously published mechanical and dynamic mechanical 

results. Furthermore, the application of this theory to these seemingly nonideal networks 

resulted in network parameters that were in excellent agreement with traditional 

equilibrium swelling estimates. 



Xlll 



The work was then extended by utilizing this ethylamine solution to catalyze the 
sol-gel reaction in-situ. The effect of this change in catalyst upon the oxygen diffusivity of 
the hybrids as a function of polysilicate loading was investigated using a luminescence 
based approach. While the diffusivity of the acid catalyzed gels decreased with increasing 
loading, the base catalyzed gels did not indicating that the polysilicate domains resulting 
from the base catalysis possess considerable porosity. However, the pores appear to be 
much too small for Knudsen diffusion, a commonly observed gas separation mechanism in 
porous ceramic membranes. 

To investigate the influence of polysilicate network polarity and spatial 
distribution, the sol-gel processing of the hybrids was adjusted to produce two classes of 
gels. One exhibited a more discrete polysilicate phase possessing greater network 
connectivity and reduced silanol content than the other. This was accomplished by using 
dimethylformamide in place of tetrahydrofuran as the organic solvent constituent of the 
sol. Poly(methacrylic acid)-PTMO-polysilicate semi-interpenetrating polymer networks 
(SIPNs) were then produced by y polymerizing monomer swollen hybrids using a 60 Co 
source. Fourier transform infrared spectroscopy was used to confirm the reduced silanol 
content. Tensile testing revealed that the better developed, less polar inorganic network 
containing SIPNs exhibit decreased elongation at failure and post-yield elongation to 
strain hardening. Thus, the origin of the exceptionally high elongation of these SIPNs is a 
diffuse polysilicate network capable of extensive hydrogen bonding and deformation under 
load. 



xiv 



CHAPTER 1 
INTRODUCTION 

This dissertation is dedicated to the study of several aspects of hybrid organic- 
inorganic composites. The pervading theme throughout the work is an attempt to discern 
the role that the polysilicate reinforcing phase of a micro-phase separated, polymer matrix 
composite plays in dictating the mechanical, thermo-mechanical and gas transport 
properties of the overall materials system. The sol-gel process is utilized for composite 
consolidation. However, as there are many variables active in this process, the work 
specifically concentrates on the effects of spatial distribution and degree of network 
connectivity of the polysilicate phase as controlled by pre- and post-gelation exposure to 
network modifying solutions. More details regarding the specific studies will be given in 
the following paragraphs as each facet of this research is developed. 

Composites play an ever increasing role in our daily lives. Driving forces such as 
reduction in weight, increases in stiffness or hardness and the ability to tailor a material 
response in a specific direction, if desirable, have spurred composite development. A 
general definition of a composite is that of a multiphase system in which one phase resides 
within the other and is responsible for the augmentation of an engineering property for the 
overall system. Such a material system is in contrast to a material blend in which one 
phase simply reduces the amount of the costlier phase while providing little to no increase 
in any engineering property. 



Many types of composites exist within the ceramic, metal and polymer industries. 
Narrowing the discussion to the realm of polymer matrix composites, a plethora of 
composite systems still exist. For example, theromosets such as epoxy, polyester and 
polyurethane have all been used as matrix materials, as have thermoplastics such as 
poly( ether ether ketone), polyimides and polyamides. Similarly, a variety of reinforcing 
fibers exist. These may be ceramic, metal or polymer. The only design criterion that 
exists is that the modulus of the fiber must be much greater than the modulus of the 
matrix. Thus ensuring reinforcement of the matrix. 

Traditionally, composite systems are classified according to the length of the fiber, 
i.e., continuous, discontinuous or particulate. However, with the advent of new 
processing techniques developed by materials chemists over the past several decades, 
composites can now be produced where the size of the reinforcing phase is on the 
nanometer scale. The development of these nanocomposites has been driven by groups 
interested in such aspects of materials science as the quantum confinement of electrons, 
self-assembling devices, high surface area catalyst supports and, an area of particular 
interest to this author, the extent that these nanophases are responsible for overall 
composite properties such as elasticity and gas transport. 

Returning to the arena of nanocomposite processing, a variety of techniques have 
been developed enabling the synthesis of nanometer-in-size reinforcing phases. Examples 
such as sol-gel-processing, micellular formation, and in-situ growth of nanoclusters 
dominate the literature. However, regardless of the specific route utilized, all of these 
techniques rely on the same underlying approach, that of materials chemistry. 



One class of nanocomposites that has been receiving a particularly large amount of 
attention in the past decade has been that of hybrid organic-inorganic composites. Hybrid 
composites, or simply hybrids, almost exclusively employ combinations of organic 
polymers or proteins and ceramic oxides or minerals. Numerous techniques have been 
developed for producing hybrids. However, with the exception of clay based composites, 
the review of hybrid technology provided in Chapter 2 will concentrate upon only sol-gel- 
derived composites. This is done to provide the reader with a detailed review of the 
processing routes, properties and applications of hybrid systems most closely related to 
the investigations detailed in Chapters 3, 4 and 5, which utilize the sol-gel-processing of 
ethoxysilane functionalized poly(tetramethylene oxide) (PTMO) and the silica precursor 
tetraethoxysilane (TEOS). 

Intrinsic to hybrid composites, such as the PTMO-polysilicate system, is the 
influence of size, shape and continuity of the exceptionally small reinforcing polysilicate 
phase. For example, considering the significant increase in surface area relative to volume 
occurring as the particle size decreases, it should be clear that the interphase becomes 
increasingly important. In the case of sol-gel-derived composites, the spatial distribution 
of the reinforcing phase becomes particularly important given that the fractal nature of the 
inorganic structures produced via the sol-gel-process lead to extensive interaction between 
the matrix and the reinforcing phase. Probing the extent of PTMO and polysilicate 
interaction as a function of post-gelation processing is the focal point of Chapter 3, 
wherein elementary rubber-elasticity theory is employed in conjunction with dynamic 
mechanical spectroscopy as a tool for quantifying the extent of phase interaction. 



Another critical issue when considering the increasing surface area contribution is 
that of surface chemistry. For example, the polysilicate networks that result from acid 
catalysis of silicon alkoxides are essentially defect laden structures containing alkoxy and 
hydroxyl groups that have not yet undergone complete hydrolysis and condensation. 
Hence the term polysilicate is used and not silica in reference to the lack of a well 
developed, oxygen-bridging network characteristic of fused silica. These unreacted polar 
species give rise to hydrogen bonding and other molecular forces that influence composite 
properties. Chapter 4 details investigations into the role that these defects play in dictating 
the mechanical response of poly(methacrylic acid)-based, semi-interpenetrating polymer 
networks (SIPNs) formed via the y polymerization of methacrylic acid swollen PTMO- 
polysilicate hybrids. These SIPNs exhibit the exceptionally high stress and elongation at 
break characteristic of engineering thermoplastics despite the fact that poly(methacrylic 
acid) is an organic glass. 

A final issue concerning the role of the polysilicate phase in determining the 
properties of sol-gel-derived hybrid composites is that of porosity. The sol processing 
conditions, such as the use of acidic versus basic catalysis, significantly influence the 
extent and continuity of the porosity present within the inorganic domains of these hybrid 
composites. Considering that the diameter of these pores may be as large as 10 nm, it is 
possible that the transport of gases is occurring via Knudsen diffusion at pressures less 
than ca. 15 atm. This transport mechanism is in contrast to the traditional solution and 
diffusion mechanism occurring within the elastomeric PTMO phase. Therefore, the 
investigations detailed in Chapter 5 addressing the transport of oxygen within these 



PTMO-polysilicate hybrids as a function of inorganic loading and spatial distribution have 
allowed an assessment of the types of porosity present within these composites. 

Detailed conclusions regarding each of these areas of investigation are provided at 
the ends of Chapters 3, 4 and 5. However, Chapter 6 will highlight some of these critical 
findings, summarize the overall impact of this research upon the field of hybrid technology 
and polymer science in general, and detail possible future directions for the research. 



CHAPTER 2 
THE DEVELOPMENT AND APPLICATION OF HYBRID COMPOSITES 

Development and Detailing of Specific Composite Systems 

As the following pages will demonstrate, many polymers and oligomers have been 
utilized in the synthesis of hybrid organic-inorganic composites. These include classes 
such as poly(alkenes, acrylates, ethers, esters, amides, imides and dienes). Similarly, a 
variety of inorganic alkoxides have been used. These include methoxides, ethoxides, 
isopropoxides and butoxides of silicon, titanium, zirconium and aluminum. Additionally, 
phenyl, vinyl and amino substituted derivatives of these silicon alkoxides have been 
incorporated into glassy networks. Despite the wide variety of components that have been 
mixed to date, the driving force for the development of hybrid composites remains to be 
the improvement in some material property arising from the synergistic combination of the 
organic and inorganic phases. Whether it is the improved reinforcement of a rubber over a 
conventional particulate filler, increased abrasion resistance of polymer glasses, improved 
temporal stability of embedded lumiphores or increased selectivity to gas mixtures 
permeating through the bulk, the key words are improved and increased. This section will 
provide some historical perspective for the development of these hybrid composites. 
However, the main focus is upon the information gained from the individual systems 
studied. 



Poly(Dimethyl Siloxane) 

Development of organic-inorganic networks began in the early to mid-1980s using 
poly( dimethyl siloxane) (PDMS) oligomers. This polymer was chosen because of the 
close similarity between the polysiloxane chains and the expected sol-gel-derived 
polysilicate chains. Hence, it is also an ideal organic-inorganic glass. Similar synthetic 
approaches were employed independently by two groups: J.E. Mark (The University of 
Cincinnati) and G.L. Wilkes (Virginia Polytechnic Institute and State University). 
However, the research goals of each group were different. Mark's approach focused on 
improving the mechanical properties of PDMS by using the sol-gel processing of metal 
alkoxides to generate inorganic fillers (silica, titania, alumina and zirconia) within the 

PDMS network. * "8 In all cases, the hybrids generated by either the simultaneous curing 
and filling or the precipitation of the inorganic component within an existing swollen 
network (in-situ precipitation) showed significantly better reinforcement than the network 
containing fumed silica, an alternate reinforcing agent. Electron microscopy of the hybrids 
synthesized using the in-situ precipitated silica show that the inorganic phase did not 

agglomerate into large particles within the original PDMS network. 3,4 Rather, most 
"particles" fall in the range of 20 to 30 nanometers and were very finely dispersed. Small 
angle X-ray scattering experiments on these reinforced networks demonstrated that under 
strongly basic conditions and large excesses of water, uniformly dense particles 

(nonfractal) resulted* Conversely, under acidic or neutral conditions more extended, 
less-dense inorganic structures resulted. Sohoni and Mark also studied the thermal 
stability of in-situ filled PDMS networks using thermogravimetric analysis (TGA) and 



demonstrated that this type of reinforcement improved the thermo-oxidative stability of 

these materials more so than fumed-silica reinforced or ordinary PDMS.8 

Simultaneous with Mark's research was that of Wilkes' group which also used 
PDMS as the polymer. Similarly, silica was used as the reinforcing phase. These 
materials were termed ceramers to denote the contributions of both the ceramic, inorganic 

glass and the polymeric, organic glass. 9 " 11 As the order of the name implies, i.e., not 
polymics, the emphasis was placed on developing sol-gel glasses containing more 
inorganic than organic material. The idea being to engineer glasses with controllable 
flexibility/brittleness rather than the reinforcement of rubbery systems as Mark was 
examining. An overview of the laboratory's contributions, as well as a general review of 

sol-gel-derived hybrid composites, has been recently published. 1 2 

Wilkes et al. research efforts examined the effect of silanol terminated PDMS as a 
function of loading upon the mechanical properties of these sol-gel-derived glasses. The 
principlal means of characterization were mechanical tensile testing, dynamic mechanical 
spectroscopy (DMS), and small angle X-ray scattering (SAXS). Since Mark had not 
utilized DMS, Wilkes et al. provided new insights into the structure-property behavior of 
these mixed-phase glasses. These insights included analysis of the tan 5 response to 
discern information regarding polysilicate-polysiloxane interactions. A series of 
publications address the effects of varying the acid concentration, amount of water and 
tetraethoxysilane (TEOS) added as well as the molecular weight of the PDMS utilized. 9 " 
1 1 These results indicated that the extent of hydrolysis and condensation that the network 
undergoes is a function of the amount of water added. Similarly, the simultaneous 



9 

hydrolysis and condensation of all alkoxysilanes may result in preferential reactions 
between inorganic TEOS monomer and itself along with silanol-terminated PDMS and 
itself. This would result in a phase separated morphology. 

Chung et al. 13 and Hu and Mackenzie 14 have prepared rubbery ORMOSILS 
using PDMS and TEOS. An ORMOSIL is an organically modified silicate. Traditionally, 
such materials are comprised of organosilanes which, as the name implies, possess one or 
more organic pendant groups bonded to the tetrafunctional silicon atom. When reacted 
via the sol-gel process, these organic groups remain in the structure and impart flexibility 
to the resulting hybrid. Schmidt et al. have researched these materials thoroughly and a 
good review is available. 15 The ORMOSILS produced by Chung et al. differ somewhat 
from these original hybrids in that TEOS is not an organofunctional moiety. Nevertheless, 
these PDMS/TEOS materials do exhibit a high degree of rubber elasticity, which is a 
function of the sol-gel processing variables. The authors have suggested and evaluated 
different models, which in conjunction with mechanical tensile testing and electron 
microscopy, have helped determine the morphology of these materials. 

Kohjiya et al. 16 and more recently Surivet et al. 17 have also published information 
on PDMS/silica based hybrids. In particular, Surivet et al. relate their work on ceramers 
to that of Wilkes's et al. discussed above 9-11 Using differential scanning calorimetry 
(DSC), DMS and SAXS, these researchers investigated the structure-property relations of 
polyurethane-co-PDMS (PU-PDMS)/silica hybrids, as well as polybutadiene-based 
hybrids discussed shortly. Their results revealed that the PU-PDMS/silica system was a 
three phase, micro-phase separated system. This phase separation was due to 



10 

thermodynamic incompatibility between the siloxane backbone and the hard segments 
(urethane-urea linkages) used to cap the PDMS chains with multifunctional alkoxysilanes. 
Conversely, the polybutadiene based hybrids could be considered two-phase systems 
because of good miscibility between the polybutadiene backbone and urethane-urea 
linkages. 17 These early results indicated that significant gains in mechanical reinforcement 
could be achieved using hybrid technology. However, the systems produced exhibited 
micro-phase separation, a characteristic feature of virtually all hybrids produced using 
organic oligomers. 
Polv(Tetramethylene Oxide) 

It was speculated by Huang and Wilkes that the tendency of the PDMS and silica 
to phase separate was due to the immiscibility of the PDMS with the water added for 
hydrolysis and generated during condensation. 18 Therefore, new polymer systems were 
sought which had better mixing characteristics with water. Oligomers of PTMO were 
chosen for precisely this reason. However, the absence of the siloxane chain structure to 
compatiblize the organic polymer and inorganic glass was a concern. Previously, only 
silanol-terminated PDMS was all that was necessary to react with the in-situ generated 
silica. To improve the compatibility of the PTMO with silica, isocyanatopropyl- 
triethoxysilane was reacted with the oligomers. This reaction resulted in each oligomer 
possessing end groups capable of undergoing sol-gel processing. Similarly, it was 
assumed that the reactivity of the ethoxy groups on both the TEOS and the PTMO end 
groups would be the same. This would eliminate any preferential condensations as 
discussed above in the PDMS systems. With these concerns addressed, work was begun 
on the new hybrids and focused on the variables of oligomer molecular weight, 



11 

functionality, TEOS content, inorganic precursor, gel age, acid concentration and type, 
and processing temperatures and methods. 10 . 11 ' 19-31 By using SAXS, DMS, 
mechanical tensile testing, infrared spectroscopy and NMR the structure-property relations 
of this class of hybrids were investigated. In particular, SAXS still revealed the presence 
of micro-phase separation within the transparent hybrids. However, the mechanical 
properties of the PTMO based system were much better than those of the PDMS/silica 
hybrids. 

The observed benefits of near-molecular reinforcement over conventional 
reinforcement, as well as the potential of increased hardness, abrasion resistance and 
improved thermo-oxidative stability, sparked interest in the hybridization of other polymer 
systems. Although the reinforcement of rubbers continued to receive much interest, 
attention soon turned to the incorporation of inorganics into higher molar mass 
thermoplastics. The following sections detail some of the processing utilized and the 
resulting properties that could be developed in these materials. 
Other Rubber Based Systems 

Mauritz and Jones have synthesized homogeneous, translucent organic-inorganic 
"alloys" combining poly(n-butyl methacrylate) (Mn = 75,000 g/mol) with tetraethyl 
titanate (TET) and tertbutyl titanate (TBT) 3 2 The thermal stability of these networks 
was studied using thermogravimetric analysis (TGA), and it was determined that the 
thermal onset of significant network degradation could be increased by the addition of the 
titania phase. DSC revealed that crystallization of the amorphous Ti0 2 to anatase, in the 
TET case, occurred at temperatures above the degradation temperature of the polymer. 



12 

Surivet et al. studied the thermo-mechanical properties of alkoxysilane-terminated 
hydrogenated polybutadiene-co-polyurethane (PU H-PBD) macromonomers of varying 
molar mass crosslinked via the sol-gel-derived silica phase. 17 Recall that these authors 
also investigated PU-PDMS of varying molar mass. Using a point source X-ray 
apparatus, SAXS investigations were made to determine the morphology of these 
heterogeneous materials. Those data, in conjunction with DMS and DSC measurements 
allowed these investigators to develop a qualitative morphological model for this new 
ceramer system. Unlike the earlier reported model for the PU-PDMS copolymers, which 
specified a three phase system composed of the pure PDMS regions, pure polysilicate 
regions and a distinct interfacial region arising from immiscibility of the urethane-urea end 
linkages, the PU H-PBD gels were distinctly two phase. This "improvement" in the level 
of miscibility is attributable to the increased solubility of H-BPD and the urethane-urea 
linkages. 33 These results highlight the importance of solubility in hybrid technology, for it 
is the level of mixing which gives rise to the improvements observed in hybrid systems. 
Polyoxazolines 

David and Scherer have synthesized a transparent, inorganic-organic glassy 
material composed of 50/50 volume percent TEOS-derived Si0 2 and 500 kg/mole 
poly(ethyloxazoline) (PEOX). Interestingly, there was no evidence of phase separation, as 
SAXS, DSC and TEM all indicated a homogeneous material. 34 The processing route and 
precursors utilized ensured that the two components were not covalently bonded. The 
absence of two-phase behavior is unusual for hybrids generated using a polymer of such 



13 

high molar mass. However, the strong hydrogen bonding ability of this polymer appeared 
to have provided the solubility enhancement necessary to prevent phase separation. 

Additional evidence of such an occurrence is provided by Saegusa who has 
demonstrated the ability of polymers containing strong electron donor groups, such as 
poly(2-methyl-2-oxazoline) to be incorporated and well dispersed into the evolving 
inorganic glass. 35 Again, by improving the affinity of the two phases, which was 
demonstrated using FTIR, a more intimate level of mixing is achievable. One consequence 
of this increased level of mixing, that these researchers discuss and that is addressed later 
in the applications section, is the ability to create hybrids of controlled porosity via 
pyrolysis of the organic phase. 
Vinylic Based Polymers 

Fitzgerald et al. have investigated the dielectric and mechanical spectra of hybrids 
produced by mixing a polyvinyl acetate) (PVAc)/THF solution with TEOS, followed by 
the addition of hydrochloric acid. 36 Mixtures were made believed to be 0, 5, 10, 15 and 
20 wt.% Si0 2 . These composites were transparent and, not surprisingly, FTfR revealed 
hydrogen bonding between the silicate network and carbonyl units of the PVAc. 
Interestingly, no shift in the Tg of the composites from that of the pure PVAc was 
observed. Similarly, the activation energies were calculated and shown to be independent 
of Si0 2 loading. However, the breadth of the tan 6 associated Tg relaxations did increase 
with increased filler content. 

Landry et al. have also produced transparent, homogeneous hybrids using a 50/50 
wt.% PVAc and TEOS mixture and acid catalysis. 37 Just as in the preceding paragraph, 



14 

DSC and DMS indicated only a slight increase in the Tg of the hybrid with incorporation 
of silica. Also, dynamic mechanical tan 6 responses indicate a strong interaction between 
the organic and inorganic phases, which lead to well-dispersed phases and high modulus 
rubbery plateaus. The absence of a discernible increase in the Tg of these PVAc 

composites is in agreement with the finding of Fitzgerald et al. described above.-'" These 
findings lead to the conclusion that hydrogen bonding is not energetic enough and that 
covalent bonds need to be present for an increase in the Tg to occur. 

Saegusa and coworkers have utilized the hydrogen bonding affinity of the organic 

and inorganic phase to produce transparent PVP-silica hybrids. 35,38 Using atomic force 
microscopy (AFM) and BET analysis these researchers demonstrated that the composite 
possessed a very dense microstructure exhibiting little porosity, i.e., the silica domains in 
this material have much less pore volume than sol-gel-derived silica. This observation is 
important in that the section on membranology later in this chapter discusses the influence 
of porosity within the hybrids. 

Similar to the work of Saegusa, Novak et al. have also found that polymers 
possessing functional groups such as amines and pyridines are soluble in the pre-gelled sol 

solutions, especially, poly(2-vinyl-pyridine) and poly(n-vinyl pyrrolidone).-*9 However, 
unlike the work of Saegusa et al., Novak's work focuses on the synthesis of non- 
shrinking, sol-gel-derived networks utilizing acid based, acrylate monomers to 
functionalize the alkoxysilanes prior to sol-gel processing. Upon hydrolysis these 
macroalkoxysilanes release the acrylate monomers as the hydrolysis by-product. These 
groups then undergo traditional free radical polymerization simultaneous with the 



15 

polycondensation reaction of the silane. As the hydrolysis by-product is polymerizable, no 
mass loss and corresponding shrinkage is observed. 3 ^ 

Extensive work has been done at the Eastman Kodak Corporate Research Labs by 
Landry et al. towards determining the chemical nature of the organic-inorganic interaction 
and the thermo-mechanical properties that result from the non-covalently bonded phases 
interacting within acrylic glasses. 37 ' 40 - 41 Like many of their coworkers, this group found 
that the ability to produce well-dispersed networks is attributable to hydrogen bonding 
between the silanol groups and, in this case, the carbonyl units in the poly(methyl 
methacrylate) (PMMA). 37 Hence, the highly hydrated, open polysiloxane chains resulting 
from the acid-catalyzed sol-gel process were shown to form more homogeneous, 
transparent hybrids. Somewhat surprisingly, these hybrids exhibited enhanced mechanical 
properties beyond the Tg. However, the mechanical properties were also affected by 
curing time and temperature 41 It was also noted that transparency was a function of the 
temperature of the substrate the films were cast on, e.g., >30°C resulted in films exhibiting 
no macrophase-separation. Base catalyzed systems, films cast on cooler substrates and 
non-hydrogen bonding capable polymers produced cloudy, poorly dispersed glasses. 
Another interesting feature of these reports is microscopy images provided that show 
asymmetric particulate formation in some of these hybrids. 

Pope et al. have studied the properties of porous, sol-gel-derived glass 
impregnated with benzoyl peroxide (BPO) initiated PMMA. The effect of silane coupling 
agents was evaluated also 42 It was found that the density, elastic modulus, modulus of 
rupture (MOR) and compressive strength of the material decreased as the volume fraction 



16 

of PMMA increases. Conversely, the refractive index increases as the volume fraction of 
PMMA increases. Methacryloxypropyltrimethoxysilane was then used as a coupling agent 
to improve the adhesion between the silica glass and the PMMA. As expected, the result 
was an increase in the MOR as the amount of coupling agent employed increased. 

Klein and Abramoff also examined PMMA impregnated sol-gel-derived silica 

gels. 43 Long-wave UV illumination was employed in addition to benzoyl peroxide for 
PMMA polymerization. This method prohibited the degradation of the silica xerogel due 
to moisture adsorption and desorption. Overall, the material behaved more like bulk 
PMMA than bulk silica, with the exception of hardness. 

Wei et al. have reported the synthesis of poly(allyl methacrylate-co-methyl 

methacrylate) containing hybrids using group-transfer polymerization. 44 The allyl 
methacrylate homopolymers (PAMA) were then functionalized by hydrosilylation of the 
allylic segments using Speier's catalyst. When mixed with tetrafunctional metal alkoxides 
in the presence of HC1 or methanesulfonic acid and water, transparent, sol-gel-derived 
glasses were produced over a wide range of TEOS:triethoxysilyl group ratios. However, 
when copolymers of the functionalized PAMA and PMMA were used, only the 
methanesulfonic acid produced transparent hybrid glasses. Similarly, lower fractions of 
the functionalized PAMA >•©• <10%, tended to produce grainy and opaque materials. The 
lack of a well-defined Tg in the DSC traces of the polyacrylate suggested good dispersion 
of the organic and inorganic phases. Titanium alkoxides were also employed, but met 
with less success. 



17 

High Performance Thermoplastics 

The lure of improved mechanical response, thermo-mechanical response and 
increased degradation temperatures attainable with hybrid technology has spurred 
development of high performance thermoplastic hybrids. For example, polymers such as 
poly(ether ketone) and polyimide have been used in the synthesis of hybrid composites. 
As was the case for the previously discussed thermoplastics, solubility continues to drive 
hybrid morphology and response. However, steps can be taken to promote mixing. 

Considering poly(ether ketone) (PEK) first, Noell et al. have successfully produced 
transparent hybrids of varying inorganic contents using triethoxysilane end functionalized 
PEK oligomers and TEOS 45 It was demonstrated that the Tg dependence on the curing 
temperature was governable by the time-temperature transformation behavior described by 
Gillham for network development, i.e., by increasing the cure temperature the diffusion 
limited curing associated with vitrification could be avoided 46 As a result, better phase 
mixing was obtainable. However, SAXS experiments confirmed the characteristic 
presence of micro-phase separation. The importance of this work is that it demonstrated 
the ability to produce transparent, highly loaded hybrids exhibiting Tgs as high as 200°C 
from 4,000 g/mole oligomers. Additionally, this oligomer based route circumvented the 
poor solubility experienced by Landry et al. in the synthesis of PMMA and other acrylic 
hybrids. 37 . 40 ^ 

A similar oligomeric approach was used by Brennan, who has employed 
transimidization to functionalize polyimide oligomers. These reactive oligomers were 



18 

subsequently used to produce polyimide/titania hybrids. 47 This technique resulted in the 
successful synthesis of transparent materials composed of 18%, 37%, and 54% titania. 

In an alternative synthesis route, Morikawa et al. synthesized hybrid glassy 
materials using polyamic acid and TEOS 48 The resulting films consisted of non- 
covalently bonded polyimide/silica species, and the results were quite different from those 
of Brennan. Specifically, transparent, organic-inorganic networks resulted only when the 
wt.% silica was less than or equal to 8% (assuming 100% conversion to an oxygen 
bridging network). Beyond the 8% limit, only opaque films resulted. All of the wt.% to 
70 wt.% Si0 2 containing films exhibited considerable flexibility. These materials are in 
sharp contrast to the PEK hybrids that contained up to ca. 30-40 vol % silica and the PI 
hybrids of 54 wt.% titania 45 - 47 This comparison demonstrates the importance of a 
mechanism for improving compatibility, e.g., covalent or hydrogen bonding. 
Nafion® Membranes 

Nafion® membranes represent a unique and interesting area in hybrid technology. 
These membranes are perfluorosulfonate ionomers possessing a morphology composed of 
3-5 nm clusters of S0 3 X (where X is a proton or cation) terminated side chains residing 
within a semi-crystalline, tetrafluoroethylene matrix 49 Their use in the preparation of 
hybrids stems from two sources. First, their physicochemical stability and excellent ion 
transport properties have resulted in their application in electrochemical cells and as 
templates/acid catalysts for chemical reactions. 50 Secondly, the presence of these well 
dispersed, nanometer-scale, polar micro-domains is anticipated to provide an ordering 
influence for an in-situ precipitated, inorganic network. 51 



19 

The research conducted on hybrids produced from these membranes is becoming 
extensive, with Mauritz's group at the University of Southern Mississippi contributing a 

significant portion. 49-55 jh e preponderance of the literature confirms that these polar 
groups do affect the in-situ precipitation of alkoxide swollen membranes. Perhaps the best 
study illustrating the influence of this template approach to hybrid synthesis was that 

utilizing SAXS." In the report, it was demonstrated that as the silicate loading increased 
from 0% to 74% (mass/mass) the correlation distance or interdomain spacing remained 
unchanged. This is in contrast to the studies of Wilkes' group using the sol-gel processing 
of triethoxysilane end functionalizd PTMO oligomers and various metal alkoxides. For 
these systems, increasing interdomain spacings with increasing silica and titania content 
were clearly observed. 30 

Recent research on hybrid Nafion® based composites has focused on their use as 

asymmetric membranes for gas separation. 54 - 55 These application oriented studies are 
detailed under the membrane application section later in this chapter. 
Clay Based Hybrids 

Clay has received considerable attention in the synthesis of hybrid organic- 
inorganic composites. Recently, an excellent review of this field has been published by 
Giannelis providing a good overview of the properties attainable in these composites. 56 
For example, researchers in this area have demonstrated the ability to achieve excellent 
mechanical reinforcement with less inorganic than alternative glass based 
nanocomposites. 57 Therefore, less of the inorganic phase is needed and weight can be 
saved. Another inherent advantage is the self-extinguishing characteristics of this class of 



20 

materials. 5 ^ The key to this approach is the ability of organic monomers to intercalate 
into the layers of these silicate based minerals prior to polymerization initiation. 

For example, Yano et al. have produced molecularly mixed composites of 
montmorillonite clay and polyimide. These hybrids were synthesized using 
montmorillonite intercalated with the ammonium salt of dodecylamine. Polymerization 
occurred in the presence of dimethylacetamide and polyamic acid, and the resulting 
dispersion was cast onto glass plates and cured. The cured films were as transparent as 
polyimide but exhibited reduced gas permeability and a lower coefficient of thermal 
expansion (CTE) than ordinary polyimides. 58 The reduced permeability is not surprising, 
as fillers often reduce the ability of gases to permeate through polymers. Nevertheless, 
applications such as low cost, more effective barrier coatings can be realized. 

Okada et al. have produced nylon-6/montmorillonite clay hybrids via intercalation 
of 12-aminolauric acid. When mixed with s-caprolactam and polymerized at 100°C for 30 
min, a nylon/clay hybrid (NCH) was produced. TEM and X-ray diffraction of the NCHs 
confirm both the intercalation and molecular level of mixing between the two phases. The 
benefits of such materials over ordinary nylon-6 or non-molecularly mixed, clay-reinforced 
nylon-6 included increased heat distortion temperature, elastic modulus, tensile strength 
and dynamic elastic modulus throughout the -150°C to 250°C temperature range. 59 ' 60 
Additional studies by this group have also demonstrated increases in elastic modulus and 
strength, without a simultaneous drop in impact strength for a 4 wt.% Nylon-6 
nanocomposite relative to that of Nylon-6. Perhaps most impressively, the heat distortion 
temperature increased from 65°C to 145°C as a result of hybridization. 57 



21 

In addition to the improved final properties of these clay hybrids, advantages also 
exist in the processing stages. For example, this group has synthesized vulcanized rubber 
sheets of nitrile rubber and montmorillonite intercalated with Hycar ATBN (a butadiene 
acrylonitrile copolymer). 61 Although these rubber hybrids showed enhanced 
reinforcement relative to both carbon black containing and pure nitrile rubber materials, 
they were easier to process than carbon black-filled rubbers. This leads to less wear on 
the processing equipment and lower energy consumption for production. 

Applications Utilizing Hybrid Technology 

The following sections highlight the major fields of science and industry in which 
hybrid composites are being developed and applied. Although some of the sections are 
short and provide only a quick summary, others such as the areas of luminescent 
technologies and membranology provide more detail. This is owing to their relevance to 
the studies presented in Chapter 5. 

When reviewing the literature on hybrid technology, the line between hybrids that 
actually combine an organic and inorganic phase to produce a synergistic change in a 
property and researcher-proclaimed hybrids merely incorporating a dispersed organic 
within the matrix begin to blur. For the applications previewed in the following sections, 
both types of materials will be discussed. However, it will be made clear if the work under 
discussion is relying solely on a dispersion and not on a property enhancement resulting 
from the combination of a multiphase system. For example, an excellent review of 
organically doped, porous sol-gel glasses has been published recently and addresses the 
incorporation of organics such as enzymes, proteins and liquid crystals (LC) within the gel 



22 

structure.^ Both the advantages and disadvantages of sol-gel processing are discussed 
with particular emphasis placed upon immobilization procedures. 
Microelectronics Usage 

Schmidt and Wolter have developed an organically modified ceramic 
(ORMOCER) system based on the sol-gel processing of phenyl and vinyl substituted 

organosilanes and silica."-' The system can be cured either thermally or using 
photopolymerization of the vinyl groups. The properties of this coating include excellent 
electrical properties, such as low dielectric constant and high surface and bulk resistance, 
even after weathering. Additionally, the material was stable up to 260°C. 

Popall et al. have developed several coating compositions based on organically 
substituted silicon and aluminum alkoxides, e.g., vinyltrimethoxysilane and aluminum tri- 
sec-butylate. After mixing these solutions with photoinitiators, the resulting sols were spin 
coated onto various substrates and photocured to produce patterned microelectronic 
devices.64 Polymerization can be induced using either a high wattage UV light or a 
frequency doubled argon laser. 
Abrasion Resistant Coatings 

Melamine, tris(m-aminophenyl)phosphine oxide (TAPO), diethylenetriamine 
(DETA), polyethyleneimine (PEI), 4,4 diamino diphenyl sulfone (DDS), bis(3- 
aminophenoxy-4,-phenyl) phosphine oxide (BAPPO), and epoxy have all been 
functionalized by the Wilkes group using a triethoxysilane coupling agent to produce 
molecules capable of sol-gel processing 65-67 These functionalized species may or may 
not be reacted with additional metal alkoxides. Once the sol-gel process is initiated, but 



23 

before significant gelation occurs, these sols can be spin coated onto substrates where they 
are allowed to finish gelling. Additionally, curing at elevated temperatures may also be 
employed. The resulting coatings are glasses composed of organic moieties or oligomers 
covalently bonded via siloxane linkages. These coatings have been shown to significantly 
enhance the abrasion resistance of a Lexan® substrate, and do so with increasing 
effectiveness as the curing temperature increases 65 * 66 Similarly, when additional metal 
alkoxides such as titanium or zirconium are employed in conjunction with the 
functionalized moieties, significant increases in the refractive index of the coatings 
result 65 

Schmidt and Wolter have also studied the ability of ORMOCERs based on 
alumina, zirconia, titania or silica and mixtures of each to function as abrasion resistant 
coatings. OJ For example, polycarbonate, when coated with an epoxy/alumino-silicate 
system, experienced a significant reduction in the degree of hazing induced by an abrader, 
as compared to uncoated polycarbonate. 
Solid Polymer Electrolytes 

Fujita and Honda have reported the successful synthesis of a transparent solid 
polymer electrolyte (SPE) based on PEO and alkoxysilanes 68 The material possessed 
good mechanical properties and high ionic conductivity, ca. 1.8 X 10" 5 S cm" 1 @ 25°C, 
dependent upon the organic-inorganic ratio and PEO chain length. 
Crosslinking Agents 

Munteanu has published a review article addressing the use of polyfunctional 
organosilanes as crosslinking agents for a variety of polyolefins 69 Most techniques 






24 

employ a grafting initiator, usually an organic peroxide, which is mixed with the polyolefin 
prior to extrusion. Upon heating, the initiator thermally decomposes to free radicals. 
These radicals then abstract hydrogen from the polyolefin backbone and thereby promote 
grafting of the organosilane onto the chain. Subsequent hydrolysis and condensation of 
the alkoxysilanes, with or without a catalyst, after polymer processing and shaping induce 
crosslinking in the preformed product. These methods have the advantage of not 
requiring treatment at elevated temperatures, e.g., above the crystalline melting point, to 
induce the crosslinking in the final part, as is the case for peroxide induced crosslinking. 
Hence, dimensional stability is maintained throughout crosslinking. 

Traditionally, peroxide and radiation were used to induce crosslinking. However, 
numerous patents exist regarding this novel organosilane crosslinking method and 
Munteanu provides an excellent review. 69 Similarly, Munteanu also reviews the complete 
spectrum of actual and potential applications of this technology. 
Nonlinear Optical Materials 

This class of materials is established owing to the nonlinear change in optical 
response that polarizable molecules exhibit in the presence of an intense electric field. 70 
The refractive index of a material is an example of a first-order, linear response whereupon 
the electron cloud of a molecule becomes delocalized, producing a dipole. However, 
under intense electric field, e.g., a laser, linear polarization fails and higher-order responses 
are observed such as frequency doubling, termed secondary harmonic generation (SHG), 
and occasionally frequency tripling, termed third harmonic generation (THG). Although 
these higher order responses are observable in the individual molecules, special processing 
is necessary to orient materials or glasses doped with these optically active species as they 



25 

destructively interfere with one another and negate any SHG. For polymers, a common 
approach is to heat the material above the Tg, expose it to a magnetic field and then 
quench/cool back below the Tg under the influence of the field. For thermosets, such as 
sol-gel-derived hybrids doped with these compounds, the equivalent processing would be 
to delay polycondensation until the magnetic field is established and allow the sample to 
gel under the fields influence. 

Numerous researchers have synthesized transparent sol-gel-derived glasses 
containing optically active compounds that give rise to nonlinear optical activity. 71-75 An 
inherent advantage of the silica gel matrix over organic glasses is its greater ability to 
stabilize the organic dopants, i.e., prevent these poled molecules from randomizing over 
time. This can be attributed to the network structure, which can be thought of as 
possessing a very high crosslink density. To enhance the temporal stability of the active 
compounds, which is a function of the mobility, researchers either covalently react the 
molecules to the gel matrix or produce optically active polymers/oligomers with decreased 
mobility. For example, equal amounts of the prepolymer of hexakis- 
(methoxymethyl)melamine and an optically active molecule created by reacting (3- 
glycidoxypropyl)trimethoxysilane and 4-[(4'-nitrophenyl)azo]phenylamine can be 
processed using sol-gel techniques to produce a hybrid composed of a highly crosslinked 
organic melamine phase covalently bonded to a polysilicate phase. Furthermore, the 
pendent group of the alkoxysilane monomer is the optically active group. 71 ' 75 The 
combination of high crosslink density afforded by the silica and melamine network and 
covalent bonding with the molecularly dispersed, two-phase matrix dramatically increased 
the temporal stability of the SHG for the poled composite. 



26 

Wung et al. have synthesized an organic-inorganic hybrid utilizing a very ingenious 
approach. 75 By mixing a sulphonium polyelectrolyte precursor for PPV and TMOS 
together and subsequently polymerizing the organic precursor via thermolysis, a hybrid 
network was formed. The catalyst required for the polymerization of the TMOS was 
supplied as a by-product of the thermolysis. Similar to the findings of others, the polymer 
doped glass exhibited improved optical quality over that of pure PPV. Although these 
researchers view the primary application of this material as a wavequide, the sol-gel 
processing of the TMOS appeared to have resulted in the shortening of the conjugation 
length in the PPV as observed in the above studies. 
Luminescent Based Instrumentation and Devices 

The porosity present in sol-gel-derived glasses and leached glasses, such as Vycor, 
have spurred development of instrumentation that utilizes this pore network to provide the 
pathways necessary for either gases or liquids to reach chemical probes dispersed within 
the gel. By incorporating an organic phase, increased ductility for the glasses and/or 
environmental stability for the embedded molecules can be achieved. The key to these 
technologies is the response of the organic lumiphore to sensitizing agents such as metal 
ions, oxygen, ammonia, etc. Nearly always, the detection mechanism is the change in 
spectral emission wavelength or intensity. Therefore, these devices rely heavily upon 
computers to monitor and quantify the spectral emissions. 

One of the more unique sensors developed has incorporated chemically active 
proteins within the pore structure of sol-gel-derived glass. These proteins exhibit changes 
in the visible absorption spectra reflective of the chemical reactions they have undergone 
upon exposure to cyanide gas and metal containing solutions. 76 The key, of course, is the 



27 






change induced in the luminescence of the molecules upon contact with the sensitizing 
agent. 

The sensing of both atmospheric and dissolved oxygen has received considerable 
attention and current efforts rely primarily upon the dispersion of an oxygen sensitive 

lumiphore within porous glass fibers and films. 77 " 79 Additionally, silica particles have 
been coated with these lumiphores, which are then generously dispersed in a polymeric 
binder and applied to the substrate. 80 ' 81 This approach to doping/coating highly porous 
glass results in sensors capable of responding to changes in oxygen concentration within 
0.5-3 seconds. However, a concern regarding the validity of the relationship between 
luminescent intensity and molecular oxygen concentration arises for glass filled, polymer 
matrix composites and organic glasses. 82 The origins of these problems and the impact 
on the hybrid composites reported in Chapter 5 will be addressed in this later chapter. 

Although this review reveals that functioning devices can be produced from doped 
silica gels, several researchers have noted that improvements can be made in the quality 
and long-term performance of these devices by the incorporation of organosilanes. 83-85 
For example, when 3-aminopropyltriethoxysilane (3-APTS) is used in conjunction with 
TEOS, the quality and clarity of the doped silica films produced are enhanced. 
Additionally, the tendency of fluoroescein isothiocyanate (FITC), the optically active 
agent, to leach from the films is reduced. 83 > 84 The reduction in leaching arises from the 
immobilization of the FITC to the gel matrix via the formation of a thiourea linkage with 
the 3-APTS. Similar increases in film quality and emission intensity were observed upon 
the addition of 3-(trimethoxysilyl)propylmethacrylate to TEOS. 85 An excellent example, 



28 

in terms of demonstrating the improvements possible, is in the production of hybrid-based 
light emitting diodes (LEDs). 

The low cost of plastics, multitude of manufacturing techniques relative to 
ceramics and the design flexibility of polymeric structures ensure that polymeric LEDs will 
be able to compete for markets such as flat panel displays. However, a key to their 
success is improving their longevity by reducing the influence of thermal vibration and 
chain mobility. To this end, recently published research has focused on the formation of a 
poly(phenylene vinylene)-organopolysilicate interpenetrating polymer network.** 6 Thin 
films were first fabricated by spin coating sols containing a PPV thermally-activated 
precursor and alkoxides of TEOS, methyltriethoxysilane and dimethyldiethoxysilane sols 
of varying molar ratios on indium tin oxide wafers. Heat treatment of these sols at 200°C 
initiated polymerization of the PPV precursor and drove the polycondensation reaction. 
These films exhibited blue shifts with increasing connectivity of the silicate phase, as 
controlled by the ratio of the organo-substituted alkoxides. As a result, a fourfold increase 
in photoluminescent intensity at the highest polysilicate loading was measured. 
Additionally, LED devices were fabricated by coating the exposed surface of these films 
with aluminum. The results showed that the electroluminescent intensity per unit current 
increases with increasing connectivity of the glass network as well. This behavior is 
attributed to the ability of the gelling glass to shorten the conjugation length of the 
polymerizing PPV precursor by hindering the molar mass build-up. In terms of lifetime, 
the hybrid glass/PPV LEDs exhibited double the usage lifetime. However, the decay in 
intensity with increasing time remained unchanged by IPN formation. 



29 

An alternative approach to altering the emission spectra of PPV is that of 
alkoxysilane fiinctionalization of PPV oligomers prior to the sol-gel reaction. Such an 
approach has been taken recently, with the focus of the work being the influence of the 
central moiety of arylenevinylene oligomers on photoluminescence. 8 ^ Although no 
devices were produced, these investigations did probe the influence of the silica network 
upon the conformation of these fixed molar mass oligomers. Unlike the just discussed 
research, which exhibited blue shifts with increasing silica gel connectivity, this research 
revealed that overall photoluminescence intensity decreases as a result of incorporation 
into a gel as opposed to in solution. Additionally, red-shifts were observed and attributed 
to the increase in planarity induced by the surrounding gel structure. 

A study investigating the thermal stability of optically active chromophores was 
published by Schutte et al. This study details the thermo-oxidative stability of coatings 
produced by either covalently or non-covalently incorporating 2,4-dinitroaniline into an 
inorganic silicate network coated onto a sapphire substrate. 88 Although some increase in 
the thermal stability of the chromophore was observed using UV-Vis spectroscopy, the 
authors conclude that this sol-gel method of chromophore encapsulation did not provide 
any real thermal and/or oxidative protection in either the covalently or non-covalently 
bonded state. 
Air/Gas and Liquid Separation Membranes 

This area of science is an extremely diverse and economically driven area of 
research at both the academic and corporate levels. 89 - 90 This research is driven by the 
gas separation industry which is interested in the rapid separation of gases such as nitrogen 



30 

from air for inert blanketing, high hydrocarbons from natural gas and hydrogen streams, 
and organic vapors from air and nitrogen streams. 89 ' 91 Essentially, only two main 
approaches can be taken when designing a membrane for gas separation: the use of a 
highly selective organic polymer or the use of a microporous ceramic. The next two 
sections detail the strengths of each class of material and develop the transport concepts 
necessary to understand the processes likely to occur in hybrid organic-inorganic 
composites. A review of the published research on hybrid membranes is then provided in 
the third section. 
Diffusion in porous materials 

The unquestionable benefits of inorganic porous membranes include exceptional 
high temperature stability, corrosion resistance to solvents, excellent mechanical integrity 
and high permeation rates. However, numerous disadvantages also exist for this class of 
membranes. For example, although a separation mechanism exists, via Knudsen diffusion, 
the selectivity of this process is less than that exhibited by polymers. This is particularly 
true of organic glasses which exhibit significantly greater selectivity and will be discussed 
in the next section. 

Considering porous materials, such as sol-gel-derived gas separation membranes or 
ultrafiltration membranes, several regimes of gas transport exist. The specific mechanism 
is dependent upon the Knudsen number, K n , which is defined as 



K„ = 



gas 



d (2-1) 



pore 



where d^ is the diameter of the pore and X $as is the mean free path of the gas. 92 
Schematically, these regimes are illustrated in Figure 2.1. If K„ is much less than 1, which 



31 



X 

ZJ 



> 

it 

b 




■Vi v^gas'^pore/ 



00 



Figure 2. 1 Schematic illustration of the relative change in diffusive flux accompanying a 
change in pore diameter for a gas of constant mean free path. 



32 

is the case for large pores relative to the mean free path of the gas, then molecule- 
molecule collisions predominate. However, the few molecule-wall collisions that do occur 
result in a loss of momentum for the molecules. Upon reflection from the wall, these 
molecules immediately collide with other gas molecules and regain their momentum. The 
continual loss and regaining of momentum via collisions leads to formation of a stable 
boundary layer moving at a speed dictated by the rate of the collisions. Momentum 
transfer ensures that all of the gas molecules in the stream move at the same pace. An 
important consequence of this process is that no mechanism of gas separation exists. 
As the diameter of the pore decreases, and X. gas begins to approach the dimensions of the 
pores, K n approaches unity. In this regime, some gas molecules collide with one another 
while others collide with the pore walls. As in the previous case, those molecules colliding 
with the wall loose momentum. However, the decreased pore diameter ensures that some 
of the rebounding molecules strike another wall before colliding with other gas molecules. 
As a result, these molecules do not regain their momentum, and the previously stable 
boundary layer breaks down. As a consequence, the speed of the gas stream begins to 
become independent of the rate of collisions. This lapse, or slip, in the momentum transfer 
gives rise to this transition regime. Continuing, as the diameter of the pore decreases 
further, virtually all of the collisions occurring are between the individual gas molecules 
and the pore walls. As a result, a complete loss of momentum transfer occurs and each 
molecule begins to travel at a velocity independent of other molecules. The speed of 
travel is dictated solely by the molar mass and collision diameter of the species. 
Consequently, gas molecules of differing molar masses, e.g., oxygen and nitrogen, would 
move at differing velocities, thereby giving rise to a separation mechanism. This regime is 



33 

termed the Knudsen diffusion regime. The limiting case of diffusion is when the Knudsen 
number approaches infinity. At this point, the diameter of the pore approaches the 
diameter of the molecule and spatial configurations of both the pore and gas molecules 
become significant, hence, the term configurational diffusion. 

As oxygen diffusion in hybrids will be studied in Chapter 5, it is useful to 
determine its mean free path at temperatures and pressures relevant to the previously 
described potential applications. Employing the kinetic theory of gases, the mean free 
path of a gas can be calculated from the expression 



i T V* RT 

gas pllMJ 



i 



(2.2) 



where p is the pressure, R is the gas constant, T is absolute temperature and M is the 
molar mass of the gas. The term r|, the viscosity of the gas, can be calculated using the 
expression 



2 
Mf8RTV' 



Tl = — 

1 2 



y-nMJ 



4iiic 2 h) 



(2.3) 



where the terms a and L refer to the molecular diameter of the gas and Avogadros 
number, respectively. Review of the above relations reveals that pressure and temperature 
are the two parameters that dictate the mean free path independent of gas type. For 
example, Figure 2.2 illustrates the effect of pressure upon the mean free path of oxygen at 
0°C and 100°C, which are the extreme temperatures most likely encountered by 
applications such as oxygen sensing instrumentation for use in wastewater treatment, 
biomedical applications and aerodynamic testing. 78,81,93 j ne upper ij m j t f 20 atm 
represents the commonly observed pressure limit utilized for research on organic polymer 



34 












100 



E 




c^ 


80 


<< 




s 




O) 


60 


X 




O 




sz 




*-> 




(0 




Q_ 


40 


$ 








C 




(0 


20 












J 1 



8 12 

Pressure (atm) 



16 



20 



Figure 2.2 The pressure dependence of the mean free path of oxygen at 0°C and 100°C, 
which are the temperature extremes most likely to be encountered in oxygen sensing 
applications. 



35 

membranes. 94,95 These calculated values are in agreement with those tabulated in the 
literature stating that the mean free paths of common non-condensible gases at ambient 
conditions fall between 80 nm and 130 nm.96 

Although strongly dependent upon processing history, the average diameter of the 
pores present in sol-gel-derived silica ranges from 2 nm to 20 nm. For example, Fosmoe 
and Hench report that Type VI gel silica is composed of interconnected pores of diameter 
equal to 10 nm 97 > 98 Therefore, Knudsen diffusion should dominate the flow, and indeed 
these researchers results confirm this 98 Similar pore diameters and Knudsen diffusion 
dependence were reported by Klein and Giszpenc for nitric acid catalyzed TEOS 

membranes." Lastly, the Saegusa group has reported BET analysis results of poly( vinyl 
pyrrolidone)/silica hybrid composites showing hydraulic radius values of 2 nm. However, 
no gas permeability measurements were reported or were indicated to be underway. 38 
Therefore, in keeping with these results, the porosity likely present within the inorganic 
domains of the hybrid composites to be produced in this study should lie in the 1 to 10 nm 
diameter range. Using 10 nm as the upper limit for pore diameter, Figure 2.2 reveals that 
at 0°C the K„ begins to approach unity at a pressure of 14 atmospheres. Given that the 
mean free path increases with decreasing pressures, the diffusion of gas through these 
pores at anticipated use temperatures and pressures will likely occur via a Knudsen 
diffusion mechanism. 

One issue remains, however, and that is the size of these pores relative to the 
collision diameter of the gases. In the case of oxygen the effective molecular diameter is 
0.346 nm. 89 If the pores within the inorganic domains of the hybrids are ca. 1 nm then it 



36 

is quite possible that the diffusion may be occurring in the configurational regime. 
Evidence supporting this speculation exists in the literature. For example, the diffusivity 
of methane in a single zeolite possessing a minimum pore diameter of ca. 0.56 nm is 8x10"* 
cm 2 /sec. Furthermore, the authors allude to the influence of pore shape when discussing 
the results. 100 jh e diffusivity of methane within a porous solid containing pores of 0.56 
nm diameter can be calculated from the relation given by Cunningham and Williams: 

r. 2 - 1 

D K =-jrv- (2.4) 

where r is the pore radius.92 Additionally, v is the average speed of the penetrant gas 
and £, is a shape factor defined by Knudsen as the diffuse reflection. Assuming that the 
diffuse reflection is equal to 1 and knowing that the average speed of a gas is given as 

- (8RTV 

v= UkJ < 2 - 5 > 

where R is the gas constant, T is the absolute temperature and M is the molecular weight 

of the gas species, then the diffusion coefficient for the Knudsen regime can be written as 

I 
4rf2RTV 



The molecular weight of methane is 16 g/mole. Given a pore radius of 0.25 nm, the 
theoretical diffusivity at 25°C is lxlO" 3 cm 2 /sec. This value is two orders of magnitude 
larger than the measured value, which was shown to be in good agreement with other 
reported values of zeolite diffiisivities. 100 These data suggest, along with other published 
results on pyrolyzed organic-inorganic hybrid membranes to be presented shortly, that 
Knudsen diffusion is not operative in materials with pores this small. Consequently, it may 



37 

not be operative in unpyrolyzed hybrids unless special precautions are taken to ensure that 
the pore structures are on the order of 5-10 nm. 
Diffusion in polymeric materials 

An excellent historical perspective of the transport of gases in polymer membranes 
was published by Stannett in 1978. 101 More recent accounts are also available detailing 
the history of the field with additional information regarding industrial interest. 89 Both 
reveal that the earliest study addressing gas transport through a membrane was conducted 
by Graham in 1829. Additionally, research in the field of membranology was well 
established by 1900 and many of the most important concepts developed by 1950. 
Research within the past two decades has focused on the use on glassy polymers owing to 
their increased selectivity. 

Gas transport through nonporous materials such as rubbers and organic glasses 
essentially relies upon the solution of the gas in the up-stream surface, diffusion through 
the bulk, and evaporation from the down-stream side of the material. 102 Credit for this 
concept is given to Graham, and the theory underlies all diffusion in polymeric 
materials. 102 > 103 However, theories differ in their treatment of diffusion through rubbers 
and organic glasses. A rubber, more specifically a polymer with a Tg less than room 
temperature, is the simplest case. Diffusion of a non-condensible gas through an elastomer 
is described by a combination of Fick's first law, 



3 " -E \dxJ ( 2 - 7 ) 

where J is the flux, D is the diffusivity, C is the concentration and x is the one-dimensional 
position of the penetrant within the bulk. Henry's law states that 



38 

c " Sp, (2.8) 

where, again, C is the concentration, S is the solubility and p is the pressure of the 
penetrant "above" the bulk. Under conditions of steady-state diffusion Fick's first law can 
be integrated from the up-stream face to the down-stream face to produce the relation 

J " D l^~"J (2-9) 

where the subscripts d and u denote downstream and upstream concentrations, 
respectively, and t is the thickness of the rubber. Henry's law can also be written with 
similar face-dependent concentrations: 

C u = Sp u and C d = Sp d . (2.10 and 2.11) 

Substitution of these Henry's law concentrations into Equation 2.9 yields 

J h = D H S(^ " t J ■ (2.12) 

The product DS is termed the permeability, P, and the subscript H denotes the Henry's 
law dependence of the flux and diffusivity. The most common method of measuring the 
permeability, diffusivity and solubility (also termed sorption) is the time-lag method. The 
enabling concept is the solution to Fick's second law for a plane sheet where the upstream 
surface at x=0 is kept at constant concentration, Ci, and the downstream surface at x=l is 
kept at constant concentration, C 2 =0, i.e., vacuum; 1 &* 

Q = TT-©J- (2-13) 

In this equation Q is the total volume of material that has passed through the sheet, t is the 
elapsed time, and 1 is the thickness of the polymer sheet. During the first few moments of 
the process adsorption is occurring and the volume of material traversing the sheet 



39 

continues to grow until conditions of steady-state flux are achieved. This established, the 
volume of penetrant having traversed the sheet becomes linear with time, and the slope is 
the permeability. Extrapolation of the linear portion of the total volume traversed, Q, as a 
function of time curve to the abscissa, i.e., solving for t when Q equals zero, produces the 
relation 

0= oTJ < 214 ) 

where is the extrapolated time value. From this equation, D can be determined. 
Knowing P and D then the average solubility, S, can be calculated. It is generally 
recognized that a time of 30 should be allowed for steady-state permeation to develop 
before the linear portion should be extrapolated and D calculated. 89 For unfilled, i.e., 
non-reinforced, rubbers this technique works very well. However, some complications are 
observed when the time-lag approach to gas transport is used on filled rubbers and organic 
glasses. 

The origin of the problem is easiest to understand in rubbers containing absorptive 
fillers such as zeolites. 105 " 107 The presence of these porous gas sinks leads to an 
increase in the time required for steady-state permeation to occur. Therefore, is 
unusually large and the corresponding gas diffusivities calculated from it, in turn, are 
unusually low. Van Amerongen accounted for the presence of these fillers by introducing 
a tortuosity factor, t, to account for the reduced diflusivity, D*, where 

D - T (2.15) 



40 

In this equation, D is the diffiisivity of the unfilled rubber. 1° 8 The term t was interpreted 
as a geometric obstruction factor introduced by the particulate. There is validity to this 
geometric origin, as 23 years later studies on block copolymers of polystyrene and 
polybutadiene of varying morphologies demonstrated geometric dependent changes in the 

diffiisivity of C0 2 . 109 However, the series of publications by Paul et al. investigating the 
absorptive nature of embedded zeolites suggested that the ability of the filler to immobilize 
the sorbed gas directly effects the value of 0.105-107 jh e equation relating and D 
(Equation 2.14) for polymers containing adsorptive fillers had been derived earlier, and 
was found to be applicable to these zeolite embedded polymers: 1 1° 

e=~[i+K%)]. ( 2. 16 ) 

In this equation the new quantity K%) incorporated a Langmuir isotherm of the form 



C 



C A bp 
A = l7^p"' ( 217 ) 



where C A is the concentration of the permeant in the filler, b is the hole affinity constant 
and C A is the hole saturation constant. 89 The results of new studies incorporating this 
dual-sorption theory showed that although the time to reach equilibrium permeation was 
increased significantly the actual change in diffiisivity was not nearly as severe as van 
Amerongen predicted. 106 

Paul's and other researchers' interest in polymers soon turned to organic glasses, 
which exhibited overall lower permeability than rubbers but significantly improved 
selectivity, that is, the ability of the polymer to transport one component of a gas mixture 
at a faster rate 95 '! 1 1-1 15 The adsorptive nature observed in the filled rubbers was also 



41 

observable in the organic glasses studied, and a postulate was soon put forth suggesting 
that since a glass is not an equilibrium structure then regions within the glass do possess 
localized order. A complete discussion of the impacts of this order upon the transport 
characteristics of glasses is provided by Kesting and Fritzsche. 89 However, at this point it 
is sufficient to say that the dual-sorption theory provides the underpinnings of the study of 
glassy polymer membranes and still finds near universal application in some form or 
another today. Furthermore, even with the complexities present in the study of organic 
glasses, the promise of increased selectivity, higher temperature separation processes and 
better mechanical strength continues to drive the development of the next generation of 
gas separation membranes. 90 

This review of the transport phenomena present in both porous and polymeric 
materials provides sufficient understanding of the multiple diffusion mechanisms possibly 
occurring in hybrid organic-inorganic composite membranes. Therefore, the following 
section reviews the research published to date on this newly emerging field of 
membranology. 
Hybrid organic-inorganic composite membranes 

The bulk of the work on hybrid gas separation membranes focuses on the use of 
the organic phase as a template for creating the desired pore size within the membrane. 
The most common approach is to blend organoalkoxysilanes with tetraalkoxysilanes to 
produce organically doped glass networks. 116 "1 19 The permeation properties of the 
membranes produced are strongly dependent upon the post gelation processing, and the 
work of Shelekhin et al. provides an excellent review relating the changes in transport 



42 

mechanisms and rates with morphological changes and pore development. 1 *6 j^ e 
majority of researchers have found average pore sizes on the order of 0.5 nm to 2 nm. 
Additionally, an absence of Knudsen diffusion was observed in all of these studies. Much 
more common were results on the permeation of a variety of inert gases, e.g., He, H 2 , 2 , 
CR,, C0 2 and SF 6 that demonstrated little dependence of He permeability with 
processing/firing temperatures. Conversely, considerable changes in the permeability of 
"larger" gases were observed with firing. This indicated that the size of the molecule was 
dictating the overall permeability. 1 16 > ] 17 As all of these gases lie in the 0.26 nm to 0.38 
nm range, the conclusion of sub-nanometer pore diameters was substantiated. 

Fortunately, several studies investigating the influence of the organic dopant levels 
upon the membranes without the use of pyrolysis do exist. 119 ' 120 An investigation by 
Smaihi et al. utilizing DSC, NMR and SAXS techniques to compliment the permeability 
data obtained on phenyltrimethoxysilane versus diphenyldimethoxysilane doped 
tetramethoxysilane revealed that the mono-substituted organoalkoxysilane gels possessed 
greater connectivity and increased He/N 2 selectivity with increasing phenyl 
concentration. 119 Conversely, the network connectivity and selectivity remained 
unchanged with increasing (di)phenyl content. Furthermore, the selectivity for these 
diphenyl derived membranes was equal to theoretical Knudsen diffusion selectivity values 
of 2.6, while the selectivity of the phenyl derived membranes reached values as high as ca. 
80 at an organoalkoxide to tetramethoxysilane molar ratio of 0.8. These results, although 
lacking in an estimation of pore size, strongly support the argument that highly 



43 

homogeneous, well developed polysilicate phases within hybrid composites/membranes do 
not exhibit Knudsen diffusion. 

In a switch from the investigations of organoalkoxysilane doped glasses, Guizard 
and Lacan have researched membranes prepared from diethoxysilane and triethoxysilane 
end functionalized terephtaloyl-based molecules. *20 Analysis of the permeability 
selectivity values for 2 /N 2 obtained on membranes resulting from acid catalysis of 100% 
of the diethoxy based reactive molecules indicated that Knudsen diffusion is not occurring. 
However, selectivity values for the triethoxy based reactive molecules processed using an 
equivalent acid/alkoxy ratio suggested that Knudsen diffusion was occurring. Thus the 
membrane possessed considerable porosity in the range of 5 nm and larger. The 
selectivity results of these triethoxysilane end functionalized molecules represent the lone 
deviation from configurational diffusion within these sol-gel derived hybrid 
membranes. ^0 



CHAPTER 3 

RUBBER ELASTICITY EVALUATED USING DYNAMIC MECHANICAL 

SPECTROSCOPY AND EQUILIBRIUM SWELLING 

Relevant Background 

One issue consistently at the forefront of composite research is the effect of the 
interphase upon the mechanical and physical properties of the composite. Similar issues 
surround hybrids and to this end previous investigations have addressed modifications to 
both the organic and inorganic phases with the goal of understanding the 
structure/property relations of these near homogeneous composites. 121, 122 

With regard to modifications of the inorganic polysilicate phase, previous results 
revealed that gels derived from an acid catalyzed mixture of 60% (mass/mass) PTMO and 
40% TEOS, which were subsequently swollen in a basic 70% ethylamine in water solution 
for up to 24 hr, exhibit enhanced phase separation of the PTMO and polysilicate 
phases. 121,123 This enhancement was not accompanied by a loss of optical transparency. 
Additionally, changes in dynamic mechanical response were observed in that the thermally 
induced syneresis typically exhibited by these polysilicate crosslinked composites could be 
eliminated by aging the gels in the basic solution. These results, as well as infrared 
evidence to be presented shortly, suggest that during the first hour of exposure the 
ethylamine treatment is the chemical analog of a thermal cure in that exposure to the basic 
solution for one hour drives the condensation reaction to virtual completion. Beyond the 
first hour, however, increased phase separation is observed in the DMS data as evidenced 

44 



45 

by the onset of PTMO crystallization. Such crystallization would occur only when the 
PTMO chains are set free of significant interactions with the vitreous polysilicate chains 
and thereby gain the mobility necessary to crystallize. It has been proposed that this phase 
separation, which does not occur on the macro-optical scale, occurs via a simultaneous 
dissolution and reprecipitation process. This process is analogous to the ripening 
observed in pure silicates exposed to basic conditions. 124 In an attempt to quantify the 
changes induced by this polysilicate ripening upon the network structure of the hybrid 
composite elementary rubber elasticity theory has been applied in conjunction with 
dynamic mechanical spectroscopy. It will be shown that this somewhat unorthodox 
approach not only provides valuable qualitative insights into the interactions present within 
these near molecularly mixed composites but also that there is excellent quantitative 
agreement with values calculated from equilibrium swelling. 

Rubber elasticity theory should serve as a useful tool for quantifying the reduced 
elasticity of the PTMO chains in that interpenetration of polysilicate and PTMO chains 
should lead to entanglements or labile crosslinks that reduce the average molar mass 
between crosslinks, M c Consequently, the more interactive or mixed that the phases are; 
the lower that the average chain length should be. The exceptional linearity of the DMS 
storage modulus versus temperature plots, to be shown shortly, suggest that the data 
could be used to calculate M c based on the thermodynamic derivation of the elasticity 
relation 1 25 

G = N v kT = =^ /-> i\ 

M c K ' 



46 

where G is the shear modulus, N v is the average number of covalently bonded, elastically 
active network chains per unit volume, k is the Boltzmann constant, R is the gas constant, 
T is absolute temperature, p is the density, and M c is the average molar mass between 
crosslink junctions. If the sample is assumed to have a Poisson ratio, u, of 0.5 throughout 
the temperature range of interest, then the elastic modulus, E, which is related to the shear 
modulus by the equation 
E = 2G(l + u), (32) 

is equal to 3G. Additionally, if the characteristic relaxation time of the gel is less than the 
time-scale of the dynamic modulus measurement, then the storage modulus, E', is 
effectively an equilibrium measurement equivalent to the equilibrium elastic modulus, E. 
Equation 3. 1 can then be updated to the form 

E' = 3N v kT = ^S. „,v 

M c (J3) 

Examination of this equation reveals that a linear regression through a plot of E' vs. 3kT 
yields a slope of N v and a crosslink density of N v /2, assuming tetrafunctional crosslink 
junctions. Equating the second and third expressions of the equalities in Equation 3.3 
produces the relation 

M PL 

N v ( 3 - 4 ) 

where L is the Avogadro constant, and again, p is the density of the gel. 

A check on the values obtained using Equation 3.4 above can be made using the 
more traditional equilibrium swelling technique in conjunction with the Flory-Rehner 
Equation for a perfect network 



47 

-ln[(l-v 2m ) + v 2m+Xu v 2 2m ] = = R (v^-^f-), (3.5) 

L J Mc * 

where V2m is the volume fraction of polymer in the equilibrium swollen mass, Vi is the 
molar volume of the solvent, and Xi,2 is the Flory-Huggins interaction parameter. 126,127 
The volume fraction of polymer in the swollen mass can be calculated using the equation 

M2 
V2 _p J 

V2m_ Vi + V2~M2 Mi < 3 - 6 ) 

P2 Pi 

where V, M and p are the volume, mass and density, respectively, and the subscripts 1 
and 2 correspond to the solvent and polymer, respectively. Note that the values estimated 
using this swelling technique have not been corrected for the volume fraction of 
polysilicate present, i.e., v^ is based on volume fraction of solids in the swollen mass 
determined during extraction of the gels in tetrahydrofuran, a good solvent. 

The x-parameter needed for determination of M c using swelling can be estimated 
from the difference in solubility parameters of the gel and solvent using 

Xw = RT * 8 ' " 5 ^ (37) 

where, again, the subscripts 1 and 2 denote the solvent and hybrid gel, respectively. 128 
The value of 6 2 for the benchmark 40% (mass/mass) TEOS-based gel containing 19 vol.% 
polysilicate has been previously measured by swelling pieces of the gel in solvents of 
different Hildebrand parameters with the solvent inducing the maximum swelling taken as 
the estimated value. 122 - 123 For this report, this same value was used for all the 



48 

ethylamine exposure times thereby assuming that swelling in the ethylamine solution 
produces no change in the Hildebrand parameter of the 19 vol.% polysilicate hybrid gel. 

These relations, having been developed under the stated assumptions, permit both 
a mechanical, elasticity-based and equilibrium swelling-based estimation of M c for the 
PTMO chains present within the polysilicate crosslinked gels. Therefore, it will be 
possible to compare the values obtained using the two approaches with insights gained 
into the effectiveness of DMS as a basis for such measurements. Secondly, it will be 
possible to use the measured values to quantify the extent of interaction, as evidenced by 
the change in average PTMO chain molar mass, between the organic and inorganic phases 
present in these hybrids. 

Experimental 

Details regarding the synthesis of the gels used in this investigation have been published 
previously along with complete mechanical and dynamic mechanical characterization 
results. 121 ' 123 Consequently, the experimental techniques and testing parameters therein 
described pertain to this report also. Briefly, a dissolved mixture of 60%:40% 
(mass/mass) PTMOTEOS, henceforth referred to as TEOS(40), were cast into 
polystyrene petri dishes from an acidified isopropyl alcohol-tetrahydrofuran (4:1) co- 
solvent system. The PTMO oligomers used in this study had been functionalized by 
reacting isocyanatopropyltriethoxysilane with 2,000 g/mole poly(tetramethylene ether) 
glycol in bulk at 70°C for 4 days to produce reactive oligomers capable of undergoing sol- 
gel processing in the presence of additional metal alkoxides, such as TEOS. This reaction 
and the resulting end functionalized oligomer are shown in Figure 3.1. After casting, the 



49 



HHHH HHHH HHHH 9 C 2 H5 

HO-^~9-c-(|:{o-c-c-(p-c)o-c-c-6-c-OH 0=C=N-C 3 H 6 — £i-OC 2 H 5 

HHHH HHHH" HHHH I u 

UU2H5 



? C2Hs HO HHHH OH OC 2 H 5 

H 5 C 2 0-Si-C3H 6 -N-C(0-9-p-6-9}0-C-N-C3H 6 -Si— OC 2 H 5 
6C 2 H 5 HHHH" ^ C2Hg 



Figure 3.1 Reaction schematic illustrating the synthesis of end functionalized 
po!y(tetramethylene oxide) via the reaction of a 2% molar excess of isocyanatopropyl- 
triethoxysilane and 2,000 g/mole poly(tetramethylene ether) glycol. 



50 

gels were covered and allowed to gel for 4 days. Subsequently, the gels were uncovered 
for 2 days to allow evaporation of any residual alcohol, water or tetrahydrofuran (THF). 
Next, the gels were immersed in THF and were allowed to swell at ambient for 24 hr. 
This swelling was followed in series by vacuum drying at 40°C and 10 Torr for 24 hr, 
swelling again in water at ambient for 24 hr and a final vacuum drying at 40°C and 10 
Torr for 24 hr. After this processing the gels were considered benchmarks and were ready 
for modification using the 70% ethylamine in water solution. 

The standardized PTMO-polysilicate gels were placed into Pyrex petri dishes pre- 
filled with the aqueous ethylamine solution (pH of 12.5) for 1, 4, 7, 13 and 25 hr. After 
swelling, the gels were removed and de-swollen in water with multiple rinses until the pH 
of the de-swelling water had returned to its proper value of 6.5 (ca. 8 hr). Vacuum drying 
under the conditions described in the previous paragraph completed the ethylamine 
processing. 

For the sake of comparison, gels were made by crosslinking the PTMO chains 
without adding TEOS. These gels are referred to as TEOS(0) gels. The sol-gel 
processing of these reactive oligomers in the absence of TEOS produced an elastomeric 
gel containing ca. 4.5 vol.% polysilicate. This gel underwent the same swelling/extraction 
and vacuum drying process as the TEOS(40) gels. 

Multiple characterization techniques were utilized for this study. Dynamic 
mechanical spectroscopy (DMS) was performed using a Seiko DMS 200(FT) interfaced 
with a Seiko Rheostation Model SDM/5600H. Testing for all composites was carried out 
from -150°C to 200°C at a heating rate of 0.75°C/min in a dry nitrogen atmosphere 



51 

maintained at an approximate flow rate of 200 ml/min. The test frequencies ranged from 
0.1 to 10 Hz with a strain amplitude of 0.1%. 

Thermogravimetric analysis was performed using a Seiko TG/DTA 320 interfaced 
with the same Seiko Rheostation Model SDM/5600H as the DMS 200(FT). A heating 
rate of 10°C/min in a dry air atmosphere maintained at a flow rate of approximately 100 
ml/min was used. The temperature range investigated was 25°C to 1,000°C. 

Swelling measurements were obtained by immersing three 9.5 mm diameter disks, 
which had been punched from the cast films using a No. 6 cork-borer, into small petri 
dishes filled with THF. The averages and standard deviations given throughout this work 
are based upon the swelling values for these three samples. 

Density measurements were obtained using a Mettler density determination kit 
employing distilled water at ambient conditions. Five samples were used for each 
ethylamine exposure time with the averages and standard deviations displayed in all 
figures. 

Fourier transform infrared spectroscopy (FTIR) was performed using a Nicolet 
20SXB FT-IR spectrometer. A Perkin-Elmer attenuated total reflection (ATR) stage was 
set to 45° using a KRS-5 trapezoidal crystal obtained from Spectra-Tech, Inc. In all 
cases, 32 scans were sufficient to collect reproducible spectra with an instrumental 
resolution of 4 cm" 1 . All spectral subtractions were performed automatically using the 
OMNIC FT-IR software package supplied by Nicolet. 



52 
Results and Discussion 

Physical Characteristics 

The density and percent residue on ignition of the benchmark TEOS(O) and TEOS 
(40) gels as well as the TEOS(40) gels exposed to the ethylamine and water solution for 
up to 25 hr are given in Table 3.1. Examination of this data reveals that, as expected, the 
density of the gels and the percent residue on ignition increase as the mass loading of 
TEOS increases from 0% to 40% of the initial sol. Similarly, there is another slight 
increase in the residue mass for the 40% samples exposed to the ethylamine for 1 hr. The 
density, however, decreases somewhat after the first hour of exposure. Changes such as 
these are possible if the number of organic defects, such as the hydroxyl groups known to 
be present in the polysilicate domains, are converted to oxygen bridges during the initial 
swelling stages thereby producing a more highly connected polysilicate phase containing 
fewer volatiles. 1 ^ a. strong base such as the ethylamine solution employed in this study 
would be an effective catalyst for such rapid refinements in the polysilicate phase and 
FTIR results presented next confirm these changes. However, for the gel density to 
decrease upon exposure to the ethylamine solution a volume expansion within the PTMO 

Table 3.1 Room temperature densities of the TEOS(40) gels investigated. 





Density @ 25°C 


% Residue on Ignition to 1,000 °C 


Gel Type 


(g/cm 3 ) 


(10°C/min, air flow at 100 ml/min) 


TEOS(0) Benchmark 


1.028 ±0.00312 


2.5 


TEOS(40) Benchmark 


1.154 ±0.00230 


20.5, 20.7 


TEOS(40) 1 Hr Exp. 


1.146 ±0.00173 


21.2,21.1 


TEOS(40) 4 Hr. Exp. 


1.145 ±0.00300 


21.0,21.0 


TEOS(40) 13 Hr. Exp. 


1.147 ±0.00207 


20.9, 20.8 


TEOS(40) 25 Hr Exp. 


1.143 ±0.00245 


20.8,21.0 






53 



phase must occur that offsets the increased polysilicate phase density. Results presented 
in the following sections will help explain this trend. Of critical importance, however, is 
that beyond the first hour of exposure of the gels to the basic solution there appears to be 
no significant change in either gel density or mass residue on ignition. Similarly, this 
exposure does not result in any detectable loss of polysilicate in the bulk gels. 
ATR-FTIR Spectroscopy 

In addition to verifying the hypothesis of silica refinement via enhanced 
condensation of unreacted hydroxyl groups during the first hour of exposure, an analysis 
of the polysilicate structures present is necessary to confirm that beyond this first hour of 
exposure no chemical changes leading to increased crosslink density are occurring. Such 
changes would artificially increase the slope of the modulus versus temperature plots 
within the rubbery regimes of these gels. FTIR provides the insight needed to examine 
these chemical changes. However, the thickness of the hybrids, ca. 0.3 mm negates the 
use of transmission FTIR. Therefore, the ATR technique was utilized on all gels. 

The contributions of the polysilicate phase in all of these hybrids can be observed 
in two regions. Figure 3.2 displays the lower wavenumber region, or fingerprint region, 
dominated by the stretching of the ether linkages in the PTMO at 1100 cm" 1 , the 
asymmetric stretching of the Si-O-Si groups at 1050 cm" 1 and the Si-OH stretch at 955 
cm -' 130-132 Additionally, subtle silicate-based absorbances are observable in the 
symmetric stretching/bending of Si-O-Si at ca. 800 cm" 1 and what has been attributed to 
skeletal motion of 4-fold siloxane ring structures at ca. 570 cm" 1 . 130 . 131 . 133 The second, 
or high wavenumber region where silicate absorbance occurs is shown in Figure 3.3 where 
absorbed water and free Si-OH stretching is observable at ca. 3300 cm" 1 . 



54 



CO 

i 

o 
w 

< 



1.0 



0.8 



v PTMO Ether -| yAsym Si-O-Si 



8 0.6 



0.4 



0.2 



0.0 



Ethylamine Solution 
Exposure Time (hrs): 




1000 
Wavenumbers (cm-1) 



500 



Figure 3.2 Low wavenumber region of the ATR-FTIR spectra of a benchmark TEOS(0) 
gel, benchmark TEOS(40) gel and TEOS(40) gels exposed to the ethylamine water 
solution for the indicated times. 






55 



CD 

o 

c 

(0 

■e 

o 
(/) 

.D 
< 

* 



J2 

I 



0.8 



0.6 



0.4 



£ 0.2 



0.0 



Ethylamine Solution 
Exposure Time (hrs): 




3600 



3400 



3200 



3000 



Wavenumbers (cm-1) 



Figure 3.3 High wavenumber region of the ATR-FTIR spectra of a benchmark TEOS(0) 
gel, benchmark TEOS(40) gel and TEOS(40) gels exposed to the ethylamine water 
solution for the indicated times. 



56 



Considering the spectra shown in Figure 3.2, the most pronounced change induced by 
continued exposure to the ethylamine solution is that of the decreasing intensity of the 
asymmetric Si-O-Si stretching and Si-OH stretching. By taking the ratio of the VMymm (si-o- 
si> to the vc-o-c peak, shown as the lower plot in Figure 3.4, it is clear that there is a 
continued decrease in the strength of the Si-O-Si band with increasing exposure. The 
depth of penetration of an IR beam is a function of wavelength, angle of incidence and the 
index of refraction of the material and crystal. 134 Based on these variables, the outermost 
2 u.m of the sample surface, or ca. 0.7% of total thickness, is being probed at 1050 cm' 1 . 
This indicates that there is a reduction in the amount of polysilicate being detected near 
the surface of the gels after a day of exposure. The TGA data presented in Table 3.1, 
however, confirms that this reduction is negligible and that the overall amount of silica is 
virtually unchanged. 

The issue that remains is what changes are the ethylamine solution inducing in the 
polysilicate phase that remains behind? This question can be answered by subtracting out 
the contribution of the crosslinked PTMO, i.e., the benchmark TEOS(0) gel, and 
examining the spectral response of the polysilicate phase alone. Figure 3.5 displays the 
resulting spectra, which are dominated by the previously discussed asymmetric stretching 
of silica and silanol groups. However, there is another peak present at ca. 1 165 cm" 1 in all 
5 of the spectra that is not immediately discernible in Figure 3.2 but is commonly observed 
in the spectra of gel derived silica. There is, however, some ambiguity surrounding its 
origin in that this peak/shoulder has been attributed to both the 3 -fold degenerate 
stretching frequencies of Si0 4 tetrahedron and the longitudinal optic mode of the 
asymmetric Si-O-Si stretch. 131 . 132 Regardless, it is a common characteristic of silica. 



g 
« 2 

c 
0) 



-*: 

(0 -J 

1 

Q_ 



O 



57 



• v Asym Si-O-Si ' v Si-OH 

O v Asym si-O-Si ' v C-0-C (PTMO) 



O 



o 



o 



o 



5 10 15 20 25 

Exposure Time to Ethylamine Solution (hrs) 



Figure 3.4 The ratio of peak intensities for the asymmetric Si-O-Si and Si-OH stretches 
(1050/955 cm" 1 ) as well as the ether linkage of the PTMO and Si-OH stretch (1100/955 
cm- 1 ). 



58 



d> 
o 

c 

CO 

1 

o 

(0 
-Q 
< 




1400 1300 1200 1100 1000 900 

Wavenumbers (cm-1) 



800 



700 



Figure 3.5 Low wavenumber region of the ATR-FTIR spectra of the TEOS(40) gels 
exposed to the ethylamine water solution for the indicated times after subtracting out the 
benchmark TEOS(0) spectra. The spectra, therefore, are those of the polysilicate phases 
in the TEOS(40) hybrids. 






59 

The fact that all three significant peaks associated with silica are observable at the proper 
locations lends credence to the subtraction and ATR technique employed in this study. 
Unfortunately, one artifact of the subtraction does become visible as the exposure time 
increases. This artifact is the ether linkage of the PTMO at 1100 cm' 1 , which becomes 
more visible as the polysilicate at the surface of the samples begins to dissolve. 

Returning to an analysis of the nature of the polysilicate phase, a technique has 
been developed by Mauritz et al. on similar hybrid systems utilizing the ratio of v^y™ si-o- 
si/ Vsj.oh(1050 cm' / 955 cm' ) as a semi-quantitative/qualitative assessment of the degree 
of silica development present in hybrids. 50 ' 54 This technique can be employed on the gels 
in this study to examine the changes induced by the ethylamine processing. The ratio of 
these peaks as a function of the ethylamine exposure time are displayed as the upper plot 
in Figure 3.4 and reveal that at least 4 hr is needed before there is an indiscernible change 
in the number of silanol species relative to oxygen-bridged Si atoms. Alternately stated, 
the first hour of exposure does not provide sufficient time for an equilibrium silicate 
structure to develop within the hybrid. Beyond this first hour, during which much of the 
absorption leading to equilibrium solvent uptake occurs, it appears that the silicate phase 
experiences no further change in its degree of condensation. A similar trend is observed in 
Figure 3.3 in that after the first hour of exposure, the number of the silanol species 
present, as evidenced by the area under the absorbance at ca. 3,300 cm" 1 , seems to remain 
unchanged and approximates the number present in the benchmark TEOS(0) gel. 

To summarize the results of the ATR-FTIR analysis, it appears that increased 
connectivity of the polysilicate phase is occurring at the surface of the gels. Based on the 
previously discussed changes in density, percent residue on ignition and DMS results 






60 

presented in the next section, these changes are most certainly occurring throughout the 
bulk of the gel as well. Additionally, some dissolution is taking place at the surface of the 
sample. However, this amount of loss is negligible and undoubtedly has little influence on 
the bulk properties of the gels. 
DMS and Estimation of Average Molar Mass Between Crosslinks 

The dynamic mechanical storage modulus, E', of the benchmark TEOS(0) and 
TEOS(40) gels, as well as the TEOS(40) gels which have been exposed to the ethylamine 
in water solution for up to 25 hr, appear in Figure 3.6. The features of interest include the 
glass to rubber transition, Tg, at ca. -80°C for all of the gels and the presence of 
crystallization and onset of crystallization in the TEOS(0) and TEOS(40) gel exposed to 
ethylamine for 25 hr, respectively. This onset of crystallization is evidence of increased 
phase separation induced by the ripening process. At temperatures greater than the Tg 
and crystalline melting point, the rubbery regimes of the TEOS(0) gel and ethylamine 
exposed TEOS(40) gels exhibit significant linearity with increasing temperature and 
prompt the analysis developed in the Introduction Additionally, there is a region between 
ca. 30°C and 130°C where the benchmark TEOS(40) gel exhibits appreciable linearity. 
The first equality in Equation 3.3 necessitates a plot of E' vs. 3kT in the rubbery regime to 
determine the average number of elastically active network chains per unit volume. This 
data is displayed in Figure 3.7, where differences in slopes are observable. Although only 
the 0.1 Hz data is shown, Table 3.2 displays the slopes, N v , for all frequencies 
investigated. The exceptional linearity is confirmed by the regression coefficient, r 2 , for 
each gel, which ranges from 0.995 to 0.998. If tetrafunctional crosslink junctions are 



61 



I 



LU 

o 



-150 -100 



50 100 



Temperature (°C) 



10.5 




CO 

Qj 
o 



Figure 3.6 Dynamic mechanical storage modulus, E', as a function of temperature for the 
indicated gels. Note the different storage modulus range for the TEOS(0) gel, which was 
done to clearly show the crystallization observed. 



62 



15 



12 - 



£ 9 

i 

o 



LLI 



— r~ 

O 
D 
A 

V 

o 

o 



1.2 



TEOS(O) Benchmark 
TEOS(40) Benchmark 
TEOS(40) 1 Hr exp. 
TEOS(40) 4 Hr exp. 
TEOS(40)13Hrexp 
TEOS(40)25Hrexp 





swtf****** 7 




anmmmmimmmi^^^ 



HMEMMMMD 



0.1 Hz 
0.75 °C/min 



1.3 



1.4 



1.5 



1.6 



1.7 



%20 



3kT x 1 ZU (Joules) 



Figure 3.7 The rubbery regime of the gels investigated expressed in terms of thermal 
energy. The slope of each line is the number of elastically active network chains per unit 
volume, N v . 



63 



Table 3.2 Number of elastically active network chains per unit volume at each frequency 
measured using the dynamic mechanical spectrometer. N v x 10' 27 (chains/m 3 ). 









Test Frequency (Hz) 






Gel Type 


0.1 


0.5 


1 


2 


5 


10 


TEOS(O) Benchmark 


0.401 


~ 


0.394 


— 


0.388 


TEOS(40) Benchmark 


2.99 


2.90 


2.86 


2.85 


2.74 


2.65 


TEOS(40) 1 Hr Exp. 


8.84 


8.79 


8.77 


8.75 


8.67 


8.53 


TEOS(40) 4 Hr. Exp. 


8.05 


8.03 


8.00 


7.91 


7.86 


7.80 


TEOS(40)13Hr. Exp. 


6.98 


7.02 


6.99 


7.01 


6.99 


6.86 


TEOS(40) 25 Hr Exp. 


4.75 


4.77 


4.74 


4.73 


4.68 


4.60 



assumed, then the crosslink density is equal to N v /2. A review of the frequency 
dependence of N v in Table 3.2 reveals that as the test frequency increases there are fewer 
chains contributing to the elastic force. This trend will be discussed in more detail shortly. 
The average molar mass between crosslinks can be calculated using the second 
equality in Equation 3.2 and the density values given in Table 3.1. Figure 3.8 displays the 
estimated M c values as a function of frequency. Recalling that the molar mass of the 
PTMO prior to triethoxysilane functionalization is ca. 2000 g/mole, values in the range of 
1550-1600 g/mole for the TEOS(0) gels are promising. This is especially true given that 
estimations of polysilicate volume based on molar mass changes upon condensation and 
additive volumes indicate that ca. 4.5 vol % polysilicate exists within the TEOS(0) gel. 
These calculations assume a 75% conversion of ethoxysilane species to oxygen-bridging 
Si0 4 tetrahedron. As expected, as the volume of polysilicate increases to 19% for the 
TEOS(40) gels the average molar mass between crosslinks decreases to a range of 230- 
260 g/mole. Both the TEOS(0) and TEOS(40) benchmark gels, each with the potential 
for thermally induced ripening, exhibit an increase in M c with increasing frequency. 



64 



1600 

1500 
300 

5 250 

o 
E 




O 



200 



150 



100 



50 



o 
□ 

A 
V 

o 

o 



TEOS(0) Benchmark 
TEOS(40) Benchmark 
TEOS(40) 1 Hrexp. 
TEOS(40) 4 Hr exp. 
TEOS(40) 13 Hrexp. 
TEOS(40) 25 Hr exp. 



o 



<y 



o 



£ 



2=£ 



0.1 



1 



Log Frequency (Hz) 




O C 



10 



Figure 3.8 Frequency dependence of the average molar mass between crosslinks, M c , for 
the indicated gels. 



65 

However, less significant increases are observable for the TEOS(40) ethylamine exposed 
gels, which were just shown to possess chemically stable polysilicate phases. Additionally, 
the values for M c are significantly reduced for the ethylamine exposed gels, and a 
minimum value of approximately 75 g/mole is exhibited by the gel exposed for 1 hr. As 
the exposure time increases to a maximum of 25 hr the value of M c increases and reaches 
a maximum of ca. 145 g/mole. 

A reduction in the average molar mass between crosslinks is to be anticipated 
based upon the incorporation of the anelastic (from the standpoint of entropy driven 
polymer elasticity) polysilicate phase. For the sake of comparison, if only the volume 
fraction of PTMO is assumed elastic, then a rule of mixtures predicts that M c values of 
1920 g/mole and 1620 g/mole should be observed for the TEOS(0) and TEOS(40) gels, 
respectively, regardless of ethylamine exposure since the total mass percent of polysilicate 
remains unchanged for all exposure times investigated (Table 3.1). As these values are 
much greater than those predicted from Equation 3.3 and displayed in Figure 3.8, the 
employment of rubber elasticity theory reveals that the near molecular level of mixing 
occurring in these acid catalyzed gels gives rise to strong restriction of the PTMO chains 
by the polysilicate chains. Furthermore, the ripening induced by the ethylamine treatment 
does indeed result in phase sharpening, as evidenced by the increasing average chain 
length between crosslinks with increasing exposure time. This interaction-based analysis 
completely explains the mechanical, dynamic mechanical, density and swelling response 
previously observed for these gels. 121 Another question remains, however, regarding the 
accuracy of the measured values, especially considering the use of dynamic moduli data 



66 

and the possible chemical curing contributions known to be operative in the as cast (0 hr) 
and 1 hr exposed samples. To test the validity of the values calculated from Equation 3.3, 
equilibrium swelling measurements were made with the goal of using Equation 3.5 to 
predict M c values for the same gels. 
Equilibrium Swelling and Estimation of Average Molar Mass Between Crosslinks 

Equation 3.5 necessitates knowledge of the volume fraction of polymer in the 
swollen state, v^, and the Flory-Huggins interaction parameter, xu, for the given 
polymer-solvent-temperature combination used. With regard to the former values, 
equilibrium mass uptake of THF for the TEOS(40) and ethylamine exposed TEOS(40) 
gels was measured at room temperature and the volume fraction of polymer calculated 
using Equation 3.6. The results are displayed in Figure 3.9 as a function of ethylamine 
solution exposure. Consistent with results predicted from the above analysis, as the 
benchmark TEOS(40) gel possess the greatest average molar mass between crosslinks, it 
swells the most and consequently has the lowest v^ value. The gel exposed for 1 hr had 
the lowest value of M c and, therefore, swells the least and has the largest v 2m value. 
Continued exposure to the ethylamine induces phase separation and frees the previously 
restrained PTMO chains thereby increasing M c , allowing more swelling and decreasing 
V2m values. 

Determination of the Flory-Huggins interaction parameter for such a complex 
composite system necessitates an experimental estimate. Previous measurements of the 
solubility parameter for the TEOS(40) based systems have been made by swelling the gel 
in solvents of differing Hildebrand parameters and using the maximum in the gel swelling 



67 



0.62 



0.57 - 



5 0.52 - 



0.47 - 



0.42 




5 10 15 20 25 

Ethylamine Exposure Time (hrs) 



Figure 3.9 Values of the equilibrium volume fraction of polymer present in the swollen 
network, v^, for TEOS(40) gels as a function of ethylamine solution exposure time. 



68 

coefficient, 19.1 MPa" 2 , as an estimate of the Hildebrand parameter of the gels. 122 
Employing Equation 3.7 and knowing that THF has a Hildebrand parameter of 18.6 
MPa l/2 and molar volume of 8.1 lxlO" 5 nrVmole, Xu is calculated to be 0.0082. The 
exceptionally low value is not unreasonable considering the affinity the polymer should 
have for THF since PTMO is made from the ring opening polymerization of this solvent. 

A concern arises regarding the incorporation of the lattice constant of entropic 
origin that some authors have put forth requiring that an additional value of ca. 0.34 be 
added to the Xu value calculated using Equation 3.7. 128 For the sake of comparison, 
calculations of M c using swelling were performed using both a Xi,2 of 0.0082 and 0.3482. 
These results are displayed in Figure 3.10 along with the values of M c calculated using 
the 0.1 Hz and 10 Hz DMS data from Figure 3.8. The overall good agreement between 
the two techniques is very encouraging considering the ease with which the dynamic 
technique can be accomplished relative to the various swelling experiments. The choice of 
Hildebrand parameter is of obvious importance and a value neglecting the entropic-derived 
0.34 contribution appears to fit best over most of the range, the exception being the 25 hr 
exposure time to ethylamine solution 

Conclusions 

Dynamic mechanical spectroscopy was performed on polysilicate crosslinked 
PTMO based gels containing an estimated 4.5 and 19 vol.% polysilicate. Additionally, 
DMS was performed on samples of the 19% polysilicate gel, which had been exposed to a 
basic ethylamine and water solution for times ranging from to 25 hr. The data revealed 
exceptional linearity in the storage modulus versus temperature plots warranting 



69 



03U 


i 






■ ■■- 


! ! 1- 

Measurement 


300 










Techniaue 




■ 






□ 


0.1 HzDMS 


250 


_l • 









10HzDMS 



O 








A 


Swelling (x 1 2 = 0.348) 


| 200 


«*: 






V 


Swelling (x, 2 = 0.008) 


o 


\\ 










150 


1 * 


















100 


i 


i 








50 


i 


1 — i 



10 15 

Exposure Time (hrs) 



20 



25 






Figure 3.10 Comparison of the values for M c calculated using Equations 3.5, 3.6 and 3.7 
based on equilibrium swelling using two values of % 1,2, as well as values for M c obtained 
using Equations 3.4 and 3.5 for the 0. 1 Hz and 10 Hz data from the DMS. 



70 

the use of elementary rubber elasticity theory as a tool for measuring the average molar 
mass between crosslinks. FTIR analysis revealed that although some chemical change 
occurred during the first hour of exposure, discrediting the predicted values of M c , 
continued exposure resulted in equilibrium polysilicate structures. The significantly 
reduced values of M c obtained using the DMS based technique relative to those predicted 
by a rule of mixture for both the benchmark and ethylamine exposed TEOS(40) gels 
confirm extensive interaction between the vitreous polysilicate phase and elastomeric 
PTMO phase. Furthermore, the increase in M c with increasing exposure times to 
ethylamine is in good agreement with the onset of phase separation observable in the DMS 
data. Most interestingly, both the trend observed for the ethylamine exposed gels and the 
magnitude of the M c values calculated were confirmed using equilibrium swelling 
calculations in conjunction with the Flory-Rehner Equation. This excellent agreement 
suggests that elementary rubber elasticity theory is a good tool for investigating the extent 
of phase interaction occurring in these hybrid composites and that dynamic mechanical 
spectroscopy provides a valid basis for collection of the data necessary for estimation of 
network parameters. 



CHAPTER 4 

STRUCTURE/PROPERTY BEHAVIOR OF ORGANIC-INORGANIC SEMI-IPNS: 

EFFECT OF POLYSILICATE LOADING AND CO-SOLVENT SYSTEM 

Relevant Background 

Previous investigations in our laboratories have focused on the development of 
semi-interpenetrating polymer networks (SIPNs) generated by the y polymerization of 
methacrylic acid, n-vinylpyrrolidone and cyclohexylmethacrylate swollen hybrid 

composites. 122 ' 123 The most striking results showed that y polymerization of 
methacrylic acid (MAA) within a poly(tetramethylene oxide) (PTMO)/polysilicate hybrid 
gel, henceforth referred to as a PMAA-PTMO SIPN, produced an SIPN that exhibited 
exceptionally high stress and elongation at break. A significant 3,000% increase in the 
tensile elastic modulus was observed over that of the benchmark hybrid. Accompanying 
the increased modulus was the observance of yielding, a mechanism that had not been 
observed in the elastomeric benchmark hybrid prior to SIPN formation. These significant 
changes in the mechanical response of the organic-inorganic hybrids were attributed to the 
formation of a PMAA-PTMO copolymer upon exposure of the MAA swollen gel to high 
energy y radiation and/or hydrogen bonding between the less than fully developed, acid 
catalyzed polysilicate network present in the hybrid and the polar acid groups of the 
PMAA. It is likely that both mechanisms are at work given the high energy 
polymerization route employed. However, by suitable choice of the solvent system it is 
possible to gain some insights into the mechanical response of the PMAA-PTMO SIPNs 

71 



72 

as dictated by the hydrogen bonding ability and spatial distribution of the polysilicate 
phase present in the benchmark gels during y polymerization of MAA. It has been shown 
that the use of dimethylformamide (DMF), a drying control chemical additive (DCCA), in 
place of tetrahydrofuran (THF) as the organic solvent constituent in the organic 
solvent/isopropyl alcohol (IP A) co-solvent solution used to process the hybrids results in 
enhanced phase separation between the PTMO and polysilicate phases of the resulting 
gel. 30 Several relevant investigations of the role of DCCAs in the sol-gel-processing of 
pure silicon alkoxide systems have been published and provide insight into the origin of 
this solvent enhanced micro-phase separation in hybrid PTMO/polysilicate composites. 

Detailed studies by Orcel et al. utilizing formamide, perhaps the most potent 
DCCA in acid catalyzed sol systems have shown that hydrogen bonding occurs between 
the formamide molecule and silanol species of hydrolyzed tetramethoxysilane. 135,136 
This hydrogen bonding was shown to decrease the rate of hydrolysis and slightly increase 
the rate of polycondensation. This change in the rate constants led to larger sol particles 
and the subsequent formation of a gel exhibiting decreased interparticle connectivity, 
surface area and silanol concentration but increased mean pore size. A structural 
investigation was conducted in neutral conditions comparing the morphologies of several 
gels resulting from the gelation of sols based on a variety of DCCAs, including formamide 
and DMF. 137 It was concluded that the dimethyl substitution reduced the effectiveness of 
the formamide DCCA i.e., DMF was less effective at producing large structural units 
possessing greater pore diameters. Fortunately, a complimentary mechanism-based 
investigation exists detailing the influence of formamide and DMF in neutral sol 



73 

systems. 138 It also compares these DCCA dependent gels to an aqueous ammonia 
catalyzed system. This study revealed several differences in the gel structures produced by 
these two DCCAs and also some similarity between the DMF and base catalyzed gels. 
The rate of hydrolysis was concluded to decrease in those systems containing either 
formamide or DMF, in agreement with Orcel et al. However, it was suggested that the 
rate of polycondensation also decreased for the DMF containing sols. This is in contrast 
to the increased rate of polycondensation previously observed for the formamide 
containing sols. Small angle X-ray scattering, gas chromatography and nitrogen 
adsorption experiments revealed that DMF containing sols did produce large gel clusters. 
Of additional importance, however, is that a significantly greater number of oxygen 
bridges were formed in DMF-based gels relative to either formamide or ammonia 
catalyzed gels. Inherent in the increased network connectivity is a reduced silanol 
concentration. With regard to particulate size, the radius of gyration values calculated 
from the Guineir region of the SAXS profiles indicated that the DMF gels were composed 
of colloidal-like particles similar in size to that characteristic of ammonia gels. 138 By 
comparison, the addition of formamide produced colloidal particles larger than the DMF 
or ammonia gels, while uncatalyzed sols resulted in the smallest radius of gyration of all 
four systems investigated. 

As the findings of the above studies yield similar results for the effect of formamide 
in either acidic or neutral sol systems, it seems analogous that DMF would behave the 
same in either acidic or neutral solutions as well. Therefore, it is possible to summarize 
that the addition of DMF to an acidic or neutral sol system results in the formation of a 
polymeric silicate, or polysilicate, network exhibiting a high degree of network 






74 



connectivity that mimics that of a true base catalyzed system. These findings are 
interesting in that the similarities in gel structure hold even though the mechanism of 
formation is completely different: electrophilic attack by protons in acidic sols versus 
nucleophilic substitution in basic sols. 124,139 

It is anticipated that hybrid organic-inorganic composites produced using DMF 
would possess a polysilicate morphology/spatial distribution similar to that observed in 
pure polysilicate systems, i.e., a polymer-like, colloidal phase exhibiting little interaction 
with the surrounding polymer matrix. In contrast, a hybrid produced using a THF-based 
sol would possess a comparatively less interconnected and consequently more highly 
interactive polysilicate phase. In addition to the difference in morphology, it is anticipated 
that the increased degree of connectivity of the DMF-based hybrid relative to that of the 
THF-based hybrid would lead to a decreased silanol concentration. Therefore, by 
synthesizing PMAA-PTMO SIPNs from hybrids resulting from gelation of these two 
different sols a method exists to examine the combined influence of spatial distribution and 
silanol concentration on the exceptional mechanical response of the PMAA-PTMO SIPNs. 
Indeed, this approach has been undertaken in this study. 

In addition to the combined effects of hydrogen bonding and spatial distribution, 
the polysilicate loading was also considered an experimental variable and adjusted to 
provide four loadings in the benchmark PTMO-polysilicate gels used for PMAA-based 
SIPN synthesis: 4.5, 11, 19 and 33% by volume. Clearly stated, the objective of this 
study was to determine the origin of the greatly improved mechanical response of these 
polysilicate containing PMAA-PTMO SIPNs. To accomplish this objective three main 
characterization techniques were employed: DMS, static tensile testing and Fourier 



75 

•transform infrared spectroscopy (FTIR). Additionally, atomic force microscopy (AFM) 
and SAXS were utilized to probe the morphology of some of these mixed-phase hybrids 
and hybrid SIPNs. 

Experimental 

Triethoxysilane functionalized PTMO was prepared by the reaction of 2,000 
g/mole hydroxyl-terminated poly(tetramethylene ether) glycol (Polysciences, Inc.) and 
isocyanatopropyl-triethoxysilane (IPTS) (United Chemical Technologies, Inc.) Prior to 
the end functionalization reaction the PTMO was vacuum dried overnight at 40°C/10 
Torr. Additionally, the IPTS was vacuum distilled and the middle 80% used for this 
reaction. The purity of the "middle 80" was verified at 99+% using gas chromatography. 
The end functionalization reaction involving a 1:2.09 molar ratio of PTMO: IPTS was 
carried out in a glycerin bath maintained at 70±5°C. Constant mechanical stirring of the 
PTMO/IPTS solution under nitrogen was continued until FTIR indicated no further 
changes in the ratio of intensities of the diminishing isocyanate peak at 2250 cm" 1 and the 
developing carbonyl peak at 1715 cm" 1 , ca. 4 days. 

The hybrid composites were synthesized by dissolving masses of 10, 8, 6 and 4 g 
of the end functionalized PTMO each into a mixture of either 8 ml of IPA and 2 ml of 
THF, "THF-based system," or 5 ml of IPA and 5 ml of DMF, "DMF-based system." All 
solvents were used as received (Fisher Scientific-HPLC grade). Continuous stirring for 15 
minutes in covered 25 ml polypropylene Erlenmeyer flasks resulted in clear solutions to 
which 0, 2, 4 and 6 g of tetraethoxysilane (TEOS) were added, respectively. After 5 
minutes had elapsed, the stoichiometric amount of water required for the hydrolysis of 



76 

every ethoxy group was added along with 0.0165 equivalents of ION HC1. Each of the 2 
co-solvent classes comprises 4 sols containing 0, 20, 40 and 60 wt.% TEOS, henceforth 
abbreviated TEOS(0), TEOS(20), etc., relative to end functionalized PTMO. The 
contents of each beaker were then cast into four polystyrene petri dishes, which were then 
covered for 4 days allowing the sols to gel to transparent films. The covers were removed 
for 2 days to evaporate any residual alcohol, water and THF/DMF, and then covered again 
and stored under vacuum for later use. These thin films were used as the benchmarks 
against which the effects of radiation and SIPN formation are compared. 

The synthesis of PMAA-PTMO SIPNs was accomplished by cutting portions of 
the benchmark gels into rectangular pieces approximately 1 . 1 cm wide and 5 to 7 cm long. 
Typical sample thickness for the inorganic loadings evaluated were between 0. 1 mm for 
the TEOS(60) gels and 0.5 mm for the TEOS(0) gels with the difference in thickness 
attributable to increasing polymerization shrinkage with increasing TEOS loading. The 
strips were cut into characteristic shapes and the initial weights recorded so that mass 
increases of each could be monitored throughout the processing. The rectangular samples 
from each TEOS loading were placed in Pyrex petri dishes each containing 45 ml of 
distilled MAA. After 2 hours the monomer swollen strips were removed, blotted dry and 
the mass uptake of monomer recorded. Each was immediately placed into a pre-filled test 
tube containing 1 0% (mass/mass) MAA in deionized water. Water was chosen because of 
its poor ability to swell the benchmark gels. 121 However, MAA is soluble in water. 
Therefore, a 10% (mass/mass) monomer in water solution was used to reduce the 
concentration gradient and hence the driving-force for monomer desorption. 



77 

The sample containing test tubes were then purged by bubbling ultrahigh purity 
nitrogen through each tube for ca. 1 5 seconds. Additionally, test tubes containing samples 
of the benchmark gels suspended in deionized water only were also purged with nitrogen 
for 1 5 seconds. The tubes were then sealed with polyethylene caps and were positioned 
radially 4 inches from a 60 Co y radiation source for 3 hours and 18 minutes. The dose rate 
was ca. 350 rads/min for a total dose of 0.069 Mrads. This total dose is in keeping with 
earlier experiments. 123 Upon removal from the source, the MAA solutions in the test 
tubes had gelled to a transparent water swollen network, which adhered to the surfaces of 
the samples when removed from the tubes. To facilitate removal of the clinging polymer 
all of the samples were swollen in water for 24 hr with multiple washings. Vacuum drying 
at 40°C for 24 hr/10 Torr followed by another 24 hr of swelling in THF and an additional 
vacuum drying completed the processing of the polysilicate containing PMAA-PTMO 
SIPNs. These SIPNs were stored in a desiccator under vacuum until tested. 

The mechanical tensile properties of the benchmark gels, y-irradiated gels and 
PMAA-PTMO SIPNs were all evaluated using an Instron Model 1122 equipped with a 
200 lb load cell at ambient conditions (23±1°C, ~50% RH). The strain rate was 2.5 
mm/min. Dumbbell-shaped samples were cut using the Type III ASTM die described in 
ASTM test D638M-84. The grip-to-grip distance was 25 mm, which combined with the 
2.5 mm/min crosshead velocity produced a strain rate of 10%/min. Between 4 and 6 
samples were tested for the benchmark gels, y-irradiated benchmark gels and the majority 
of the PMAA-PTMO SIPNs at the 0, 20, 40 and 60% (mass/mass) TEOS loadings. 
However, only 2 samples were available for testing the TEOS(0) THF-based benchmark 
gel and TEOS(0) PMAA-PTMO THF-based SIPNs due to material limitations. Densities 



78 

were obtained using a Mettler density determination kit relying upon Archimedes' 
principle and water with between 8 and 10 samples per datum. The error bars shown on 
all figures represent one standard deviation from the mean. 

Dynamic mechanical spectroscopy was performed using a Seiko DMS 200(FT) 
interfaced with a Seiko Rheostation Model SDM5600H. Testing for all of the SIPNs was 
carried out from -145°C to 300°C at a heating rate of 0.75°C/min in a dry nitrogen 
atmosphere maintained at a minimum flow rate of 200 ml/min. All spectra presented are 
those obtained at 1 Hz. 

Fourier transform infrared spectroscopy was performed using a Nicolet 20SXB 
FT-IR spectrometer. For the poly(tetramethylene ether) glycol and end functionalized 
PTMO transmission spectroscopy was performed using NaCl crystals with 32 scans 
collected. However, for the thicker, cast gels it was necessary to utilize a Perkin-Elmer 
attenuated total reflection (ATR) stage set at 45° and a KRS-5 trapezoidal crystal 
obtained from Spectra-Tech, Inc. In the case of the ATR, 128 scans were sufficient to 
collect reproducible spectra. For both techniques, the instrumental resolution was 4 cm" 1 . 
All spectral subtractions were performed automatically using the OMNIC FTIR software 
package supplied by Nicolet. 

Atomic force microscopy was performed in tapping mode using a Nanoscope III 
instrument manufactured by Digital Instruments on the fracture surfaces of samples loaded 
while immersed in liquid nitrogen. Small angle X-ray scattering data was acquired using a 
Siemens Kratky camera employing a M. Braun position-sensitive detector from Innovative 
Technologies. 



79 
Results and Discussion 

Physical Characteristics 

To facilitate understanding of the effect of polysilicate loading upon SIPN 
formation the experimental results obtained will be presented as a function of the 
estimated volume of polysilicate present within the gels. Table 4.1 lists the densities of 
both classes of the benchmark hybrids and the y-irradiated gels as a function of TEOS 
loading. The estimated polysilicate volume percentage is based upon the benchmark gel 
density measurements and calculations similar to those of Huang et al. accomplished by 
scaling the density with the molar mass change as the sol converts to a gel. 19 A maximum 
75% conversion of the TEOS to an oxygen-bridging network of Si0 2 is assumed. 129 It is 
evident from the density measurements that neither the co-solvent system employed nor 
exposure to the y radiation results in a significant change in the gel density at the loadings 
investigated. 



Table 4. 1 Densities of the benchmark and y-irradiated gels, as well as the estimated 
volume of polysilicate based upon the benchmark PTMO-polysilicate gel densities and 
calculations similar to those of Huang et al. 19 











Estimated 




Co-solvent 


Benchmark Gel 


y-Irradiated Gel 


Volume 


TEOS(X) 


System 


Density (g/cm 3 ) 


Density (g/cm') 


Polysilicate (%) 





1:1 IPA:THF 


1.028±3.12E-3 


1.029±3.68E-3 


4.5 





4:1IPA:DMF 


1.03O±3.87E-3 


1.030±3.87E-3 


4.5 


20 


1:1 IPA:THF 


1.071±2.34E-3 


1.074±2.34E-3 


11 


20 


4:1 IPADMF 


1.076±1.80E-3 


1.077±3.97E-3 


11 


40 


1:1 IPATHF 


1.147±2.97E-3 


1.150±2.48E-3 


19 


40 


4:1 IPA:DMF 


1.149±5.45E-3 


1.142±5.24E-3 


19 


60 


1:1 IPATHF 


1.261±6.76E-3 


1.273±1.16E-3 


33 


60 


4:1 IPA:DMF 


1.254±8.97E-3 


1.254±6.83E-3 


33 



80 
FTIR Spectroscopy 

The structure and connectivity of the polysilicate phase present within the hybrid 
gels prior to swelling in MAA and y polymerization is important in terms of understanding 
the interactions that may be driving monomer uptake, morphology and ultimately, the 
mechanical response of the SIPNs. To begin, the transmission spectra of the molten, dry 
poly(tetramethylene ether) glycol and the subsequent end functionalized PTMO after 
reaction with the IPTS are shown in Figure 4.1. The primary features of the polyether 
glycol include the C-O-C stretch at 1113 cm" 1 , CH 2 bending (scissoring, twisting and 
wagging) in the 1200 cm' 1 to 1550 cm" 1 region, CH 2 rocking at 747 cm" 1 and symmetric 
and asymmetric CH 2 stretching at 2857 cm" 1 and 2940 cm" 1 , respectively. 140 The glycol 
groups are observable via the OH stretch centered at 3477 cm" 1 . 140 After reaction with 
the IPTS several changes are evident, with perhaps the most pronounced being the 
presence of a carbonyl at 1715 cm" 1 resulting from the reaction between the isocyanate 
groups of the IPTS and alcohol groups of the ether glycol to form a urethane linkage. The 
N-H stretch of this secondary amine is observable in the characteristic bimodal absorbance 
centered at ca. 3325 cm" 1 . 140 Less significant changes are also observable. For example, 
the appearance of asymmetric stretching of the Si0 3 groups on the chain ends at ca. 790 
cm" 1 , CH 3 rocking from the ethoxy groups at 960 cm" 1 , C-N stretching at 1240 cm' 1 , and 
NH bending from the secondary amine at 1560 cm" 1 . 140,141 

The end functionalized oligomers having been characterized, the next step is an 
analysis of the gels resulting from mixing of these oligomers with TEOS in the THF- and 
DMF-based solvent systems. As the changes in the polysilicate phase resulting from 



81 



1.0 



Poly(tetramethylene ether) glycol 




0.2 



End functionalized PTMO 



0.0 
4000 




jAu 



3000 2000 1000 

Wavenumbers (cm-1) 



Figure 4.1 Transmission FTIR spectra for poly(tetramethylene ether) glycol and the 
subsequent triethoxysilane functionalized poly(tetramethylene oxide). 



82 

gelation of these two sol classes are of primary importance, the ATR-FTIR results of the 
silicate "finger print" region of the benchmark gels employing each co-solvent system at 
the four loadings investigated are displayed in Figure 4.2. Consider first the 4.5% 
polysilicate loaded gel, which possesses a spectrum dominated by the IR characteristics of 
PTMO described in the preceding paragraph. Significant changes are observable in the 
500 cm' to 1250 cm" region, however, and a review of the infrared characteristics of 
silica is in order. 

The major absorbances known to occur in this region for gel-derived silica are 
given in Table 4.2 along with other relevant information. As the lower limit of the 
detector in this investigation was ca. 460 cm" 1 , it is not possible to examine the absorbance 
band occurring in this low wavenumber region. However, three major absorbances can be 
investigated: the silicate bending region at 800 cm" 1 , the silanol stretching at ca. 960 cm" 1 
and the asymmetric Si-O-Si stretching in the 1050 cm" 1 to 1220 cm" 1 region. To verify 
that the absorbances occurring in these regions of the benchmark hybrids are attributable 
to the polysilicate phase, the ratio of the maximum of each of these absorbances to the 
PTMO associated ether stretch at 1100 cm' 1 was plotted versus the estimated volume of 
polysilicate present in the gels. These results are shown in Figure 4.3 for both co-solvent 
systems, where it is evident that the intensities of each of the silicate associated 
absorbances increases nearly linearly in intensity with increasing polysilicate loading. In 
addition to the increase in intensity ratio with increasing glass loading, the 1040 cm" 1 
absorbance associated with asymmetric Si-O-Si stretching is more intense for the DMF- 
based gels. Conversely, the defect-associated absorbances at ca. 565 cm' 1 and 950 cm" 1 
are less intense for the DMF-based gels These trends, which will be substantiated in an 



83 



Volume of 
Polysilicate (%): 




1800 1600 1400 1200 1000 800 600 

Wavenumbers (cm-1) 



Figure 4.2 ATR-FTIR spectra of the silicate fingerprint region of the THF-IPA and DMF- 
IPA benchmark gels at the indicted polysilicate loadings 






84 



Table 4.2 Primary absorbances occurring in the infrared for gel derived silica. 



-K* 



Wavenumber (cm" ) 

450141 
468142 

475131 

550132 
578130 
550-600133 



Optically Active Regions and Absorbance Assignments 
Silicate Bending/Rocking 

Rocking motion of oxygen 1 to Si-O-Si plane, some Si motion 
Si-O-Si bending 
Lower wavenumber indicates increasing network rigidity 

Decreases in intensity with heat treatment (can be eliminated) 
Si-OH rocking (absorbance is not observed in fused silica) 
Skeletal motion of 4-fold siloxane rings 



800 130,132,133,140 

805142 

800-810131 



940-970130-132 
957-975 !43 

1090(1185)131,144 

1108(1190)142 

1080(1163130,14^ 

1220132) 

1000 5 1> 54 

1050 and 1080 5 !, 54 

1080 and 1155143 

1155-^1190-1210143 



Symmetric stretching of oxygen along the bisector of the Si-O- 
Si bridge with some Si motion 

Si stretching with broadening possibly due to random packing 
of Si04 units and Si-O-Si bending 

Si-OH Stretchin g 

Si-OH or Si-0 stretching. Reduced intensity with heating. 132 

Si-OH that increases in wavenumber with increasing gel age 

Asymmetric Si-O-Si Stretching 

Si0 4 tetrahedra stretching. More intense if base catalyzed. 

Si-0 stretching (spectrum similar to low temperature tridymite) 

Transverse optical component. 13°, 141 jh e shoulder is the 

longitudinal optical component, which decreases in intensity 

upon heat treatment of gels. 132 

Tentatively assigned to branched structural component 

Linear and cyclic structural components, respectively 

Observed during first 7 hr of polycondensation 

Evolution to higher wavenumber of 1155 cm" 1 absorbance as 

age approaches 500 hr 



1630131,132 



3200-3450144,145 
3650-3680130,131 



Molecular ly Bonded Water 
-Low wavenumber region- 
Molecular water vibration from hydrogen bonded water 
(disappears upon drying and/or heat treating to <700°C) 

-High wavenumber region- 
Absorbed water and/or hydrogen bonded Si-OH groups 
Free Si-OH stretching 



* Parentheses denote approximate locations of significant shoulders. 






85 




7 14 21 28 

Volume of Polysilicate (%) 



35 



Figure 4.3 The ratio of the asymmetric Si-O-Si stretch (1040 cm* 1 ), Si-OH stretch (960 
cm" 1 ) and Si-OH rocking/siloxane ring stretching (565 cm' 1 ) to the ether stretch attributed 
to the PTMO (1100 cm" 1 ) as a function of the polysilicate loading for both co-solvent 
classes. 



86 

analysis presented shortly, suggest that the polysilicate phase generated in the DMF-based 
sol possesses a better developed glass network, i.e., is less defect ridden and closer to a 
melt derived glass than the THF-based polysilicate phase. 132 ' 144 

To more closely examine the structure of the polysilicate phases present in each gel 
the 4.5% polysilicate containing gel was subtracted from each of the 3 higher loadings to 
obtain the spectrum of the "TEOS derived" glassy phase. The peaks present between 
1300 cm" and 1600 cm' 1 are of unchanging position and common intensities for all 4 sets 
of spectra shown in Figure 4.2. Therefore, they are used as the reference peaks for 
conducting the spectral subtractions that appear in Figure 4.4. The complete absence of 
the reference peaks in the subtracted spectra lend credence to their use. 146 At the low 
wavenumber end of the spectra, the skeletal motion of the siloxane rings (possibly 
augmented by Si-OH rocking) is observable at ca. 565 cm" 1 . 130,133 As the wavenumber 
increases, the symmetric Si-O-Si stretching, also assigned to Si-O-Si bending by some 
authors, can be seen at ca. 790 cm" 1 . 1 30- 132, 142 si . H stretching is observable at 950 
cm" 1 as is the asymmetric Si-O-Si stretching at 1040 cm" 151 > 54 >130-133,143,144 ^ 
distinct absorbance at 1160 cm' 1 exists, and is in contrast to the characteristic broad 
shoulder typically present in the spectra of gel derived and fused silica. Nevertheless, the 
location of this absorbance and its relative intensity to the asymmetric Si-O-Si stretch at 
1040 cm" 1 is in good agreement with the literature, which traditionally assigns this band to 
the longitudinal optical component of the stretch. 130-132, 141,142 A fi nal comment ; s 
due regarding the bands at ca. 1700 cm" 1 , which are composed of two mechanisms. The 
first is the aforementioned carbonyl at 1715 cm' 1 , while the second, centered at ca. 1680 



87 



-Q 
< 



1.0 r 



0.8 - 



<u 0.6 
o 

5 



• 0.4 



0.2 - 



0.0 



Volume of 
Polysilicate (%): 




1800 1600 



1400 1200 1000 800 
Wavenumbers (cm-1) 



600 



Figure 4.4 ATR-FTIR spectra of the polysilicate phases present in the gels synthesized 
from sols employing tetraethoxysilane. Spectra were obtained by subtracting the 4.5 
vol.% gel spectra from each of the three higher loadings. 









88 

cm" 1 , is most likely attributable to deformation of water that is hydrogen bonded to the 
silanol species present in the gels. 131 ' 132 This explains why the apparent carbonyl peak 
does not vanish upon subtraction of the 4.5% gel from the higher loaded films. 

An examination of Figure 4.4 reveals that the absorbance bands associated with the 
DMF-based gels (aside from the dominant 1040 cm' 1 to which each spectra has been 
scaled) are all less intense than the corresponding THF-based gel bands. This is in 
agreement with the comments made earlier postulating fewer defects in the DMF gels. 
Figure 4.5 displays the ratio of the bands associated with gel silica to the primary Si-O-Si 
asymmetric stretch at 1040 cm" 1 . It can be seen that for all three ratio sets, the DMF gels 
exhibit less intense gel associated absorbances. This is in agreement with several 
published investigations which have shown that the band located at ca. 550 cm" 1 for gel 
derived silica, which is only a weak shoulder in fused silica, disappears upon heat 
treatment at elevated temperatures, as does the silanol stretch at 960 cm' 1 . 132 
Additionally, a decrease in the intensity of the 1196 cm' 1 shoulder for the DMF gels 
compared to equivalently loaded THF gels provides additional evidence that the DMF gel 
possesses a network structure closer to that of fused silica. 130 The one exception to the 
claim of increased network ideality is that of the 790 cm" 1 peak, which is less intense for 
the DMF-based gels. This contrasts the typical increase in intensity observed for gel silica 
when heat treated. 132 It is not clear at this time why the intensity is not greater for the 
DMF-based gels relative to the THF-based gels. 

Overall, there appears to be a preponderance of evidence supporting the claim that 
the polysilicate phase generated in the presence of DMF possesses a vitreous structure 



89 



o 



10 

8 



1 

o 



o 
If) 



0.4 

0.2 

0.0 
0.9 

0.6 

0.3 



O 




Co-solvent System 
O THF-IPA 




CO 



0.9 r 



0.6 



0.3 




10 15 20 25 30 
Volume of Polysilicate (%) 



35 



Figure 4.5 The ratio of Si-OH rocking/siloxane ring stretch (565 cm-1), Si-OH stretch 
(960 cm-1) and high wavenumber shoulder of the primary absorbance band in silica (1 196 
cm-1) to the characteristic asymmetric Si-O-Si stretch at 1040 cm-1 as a function of 
polysilicate loading for both classes of gels. 



90 

closer to that of fused silica, i.e., containing fewer defects such as silanol groups. This 
analysis paired with the investigation described in the Introduction utilizing SAXS to 
probe the structure/spatial distribution of the gel compliment each other well and establish 

the differences in the gels produced from these two co-solvent systems. -*0 In the case of 
the THF-based gels, a highly interactive, spatially diffuse polysilicate phase containing a 
high number of silanol species exists. In contrast, the DMF-based gels are comprised of a 
less spatially diffuse polysilicate phase containing fewer silanol species. Having 
established the characteristics of the polysilicate phase present in these two classes of gels, 
the influence upon the mechanical response of PMAA-based SIPNs can be studied. 
Effect of Radiation on Mechanical Tensile Response 

Before presenting the results of the mechanical tensile testing it is necessary to 
point out a disparity that exists in the age of the THF compared to DMF benchmark gels, 
y-irradiated gels, and PMAA-PTMO SIPNs when tested. In the case of the THF class of 
gels, 29 days had elapsed after casting before testing was conducted, while 52 days had 
elapsed for the DMF class of gels. Huang et al. have shown that as the age of a THF- 
based, 50% (mass/mass) TEOS gel increases the modulus increases and the elongation to 
break decreases. 19 Specifically, gels tested after 24 days and 53 days exhibited Young's 
moduli of 2.8 and 4.8 MPa, respectively, while the elongation at break decreased from 
68% to 44% as a result of aging for an additional 29 days. SAXS studies on these gels 
revealed an increase in the scattering intensity with no change in the peak intensity 
position or tail region decay. Both the mechanical and X-ray scattering changes are 
attributable to an increased degree of network connectivity owing to continued 



91 



polycondensation of the less than fully developed polysilicate phase. These findings are in 
agreement with similar measurements made by Scherer on sol-gel-derived, TEOS-based 
gels synthesized using similar water and acid concentrations but no organic polymer. 147 
Considering that the FTIR analysis above demonstrated a more highly developed 
polysilicate network in the DMF gels, it is anticipated that the magnitude of the changes 
exhibited by aging of the gels described in this work would be reduced significantly. 
Similarly, the influence of the polysilicate aging becomes less pronounced as the loading of 
polysilicate decreases. Therefore, although the ages of the gels are given as a reminder, 
the true effect of the age difference is considered relatively insignificant when compared to 
the magnitude of the changes exhibited after formation of the PMAA-PTMO SIPNs. 

The effect of high energy y radiation upon gels synthesized using 40% (mass/mass) 
TEOS was previously shown to increase the network crosslink density thereby increasing 
the elastic modulus and decreasing the swelling in THF, a good solvent. 122 > 123 To more 
completely evaluate the influence of the glassy phase upon the radiation induced 
crosslinking of the benchmark gels the elastic modulus, stress at failure and elongation at 
failure as a function of the polysilicate loadings were determined for both co-solvent 
systems. These results are displayed in Figure 4.6. Typical engineering stress-elongation 
responses are not shown as the gels exhibit linearly increasing stress with increasing 
elongation up to the point of failure given in Figure 4.6. Alternately stated, no yielding 
was observed for any of the inorganic loadings evaluated either before or after y 
irradiation. 

Consistent with our earlier reports, the irradiation of both classes of the benchmark 
gels resulted in an increase in the elastic modulus and decrease in elongation to failure. 



92 



CO 
O CO 

"to 

J3 

LU 



0) 



400 
300 
200 ■ 
100 - 
■ 
30 



£^20 



*£ 

CO ^ 

(0 

I 

(fi 



10 



L 



1 1 1 1— 

-♦— THF/IPA, As-cast 
-o— DMF/IPA, As-cast 
••- THF/IPA, y-irradiated 
o- DMF/IPA, y-irradiated 



o 





i— 

CD 

LL „. 

*- >o 

(D £< 

b> 

c 

LU 



»u 








60 






30 
n 


- L = 2.5 cm 

10 %/min strain rate 

i i i 







5 10 15 20 25 30 

Estimated Volume of Polysilicate (%) 



35 



Figure 4.6 Effect of polysilicate loading upon the tensile mechanical response of the 
benchmark and g irradiated PTMO-silica hybrids exposed to 350 rads/min for a total dose 
of 0.069 Mrads. The age of the THF-IPA gels was 29 days, while the age of the DMF- 
IPA gels was 52 days. 



93 

However, the stress at failure decreased for both co-solvent systems upon irradiation, 
which is a reversal of previous observations and will be addressed shortly. 122,123 j ne 
magnitude of the change in modulus and elongation at failure increases as the inorganic 
loading increases, suggesting that the y radiation interacts with the polysilicate phase 
possibly leading to its densification via localized heating. This increased degree of 
condensation should manifest itself in an increased gel density. However, as the data in 
Table 4. 1 indicates, there is no apparent increase in overall density of the hybrid gels, 
seemingly discounting that minimal additional consolidation occurs as a result of 
irradiation at this dose rate. The overall magnitude of the changes for the TEOS(40) gel 
observed in this study are less than those observed in previous investigations. ^2, 123 
Although the total dose is the same, the dose rate in prior investigations was 580 rads/min. 
This increased dose rate most likely lead to more heat being generated, a higher extent of 
silicate consolidation and the increased elastic modulus and stress at failure observed 
previously. 
Monomer Swelling and Polymer Formation 

The effect of the inorganic loading upon the mass uptake of MAA and subsequent 
polymerization to PMAA as a function of the co-solvent system and polysilicate loading is 
displayed in Figure 4.7. It is evident that within each solvent system as the volume of 
polysilicate increases, the swelling of the benchmark hybrid decreases. However, even 
though less monomer is absorbed as the volume of polysilicate increases, more is retained 
within the hybrid structure after y polymerization. This behavior suggests that the PTMO 
phase absorbs the MAA but that both the monomer and subsequent polymer possess an 



94 



90 



^ 






80 


^-J 




c 




<u 




(0 

0) 


70 


i_ 




CL 




$ 


60 


^ 




CL 




T3 


50 


§ 




i 


40 


2 




<«— 




o 


30 


<D 




E 




3 
O 


20 


> 





i i 


i 


i 


i 

• ■ 


i --1 

- MAA absorb, in THF gel 

- MAA absorb, in DMF gel 
• PMAAform. in THF gel 

■ PMAA form, in DMF gel 


D 












"*■■::■ 


■ >>. 






- 




' Q- . 




"■On!" ■ 


- 











10 



5 10 15 20 25 

Estimated Volume of Polysilicate (%) 



30 



Figure 4.7 The effect of polysilicate loading upon the equilibrium MAA absorption and 
PMAA SJPN formation for the benchmark PTMO-silica hybrids. 



95 

affinity for the polysilicate phase. Furthermore, this affinity is greater in the THF-based 
gels as they absorb more monomer and retain more polymer after y polymerization for all 
polysilicate loadings investigated. However, the lower modulus undoubtedly contributes 
to increased swelling also. The swelling response of these hybrids can be explained in 
terms of two competing mechanisms: the kinetics of monomer desorption versus polymer 
formation. 

At low polysilicate loadings, and hence minimal restriction of the PTMO chains by 
the vitreous polysilicate chains, both the concentration gradient and hydrophobic nature of 
PTMO ensure that the rate of MAA desorption will be great. As the inorganic volume 
increases, the vitreous chains restrict the mobility of the PTMO chains and hinder their 
ability to collapse in the aqueous medium. Consequently, the magnitude of this 
contribution to monomer desorption is diminished. At the highest polysilicate loading 
investigated extensive interaction and restriction of the mobility of the PTMO chains 
occurs thereby resulting in the slowest rate of MAA desorption. The limiting case would 
be a 100% microporous silica gel in which the only mechanism for monomer desorption 
would be concentration induced Fickian diffusion. The reduced monomer desorption rate 
with increasing polysilicate loading allows the formation and retention of increasing 
amounts of PMAA. As a consequence of increased PMAA formation the hydrophobic 
nature of the PTMO is offset, thus lessening the rate of desorption as polymerization 
continues. In addition to the PTMO influence, polar interactions between the acid groups 
of the MAA and the hydroxyl containing, less than fully developed polysilicate networks 
diminish the driving force for desorption of MAA from the gels. The overall decrease in 
both monomer and polymer absorption for the DMF-based gels relative to the THF-based 



96 

gels, especially at the higher polysilicate loadings, is attributable to the greater elastic 
modulus of the DMF-based gels and a lower propensity for the hydrogen bonding. It 
seems reasonable that the preferred morphology resulting from the above-described 
swelling and formation kinetics would entail PMAA-PTMO SIPN formation adjacent to 
and possibly interpenetrating with the polysilicate rich domains. 
Mechanical Response of PMAA-PTMO SIPNs 

The typical engineering stress as a function of the percent elongation for each of 
the polysilicate loadings is shown in Figure 4.8 for the THF-based SIPNs and Figure 4.9 
for the DMF-based SIPNs (note the difference in axis range). The elastic tensile moduli 
and stress and elongation at failure are given in Table 4.3 for both. The elongation values 
shown and reported are calculated using the crosshead extension as a measure of 
elongation up to the yield point. Beyond the yield, the gauge length of the Type III die, 
7.5 mm, was used. This switch in gauge length accounts for the viscous flow that visibly 
occurs only in the parallel-sided test section of the dumbbell shaped sample. The result, 
however, is artificially high elongation. Preliminary results utilizing a video camera and 
tick marks 7.5 mm apart within the parallel sided test section indicate that the true 
elongation at failure are approximately half the value shown and tabulated. Nevertheless, 
the relative differences in elongation for the samples tested at the various volume 
percentages of polysilicate are valid and comparable to those reported previously for 
PMAA-PTMO SIPNs synthesized using higher dose rates of y radiation. 123 with regard 
to the mechanical response of the 1.6% polysilicate DMF gel and 1.8% THF gel, several 
of the samples tested were ca. twice the thickness of the remaining 3 loadings owing to the 
large monomer absorption and subsequent polymer formation (Figure 4.7). Experiments 



97 



50 



40 



CO 

0_ 



$> 30 
g 
CO 
O) 

I 20 
a) 

c 

£ 10 



- 



_L 



10 %/min strain rate 
L = 2.5 cm 




Estimated Volume 
of Polvsilicate (%) 

1.8 



5.5 

11 

23 



_L 



50 100 



150 200 250 
% Elongation 



300 350 400 



Figure 4.8 The effect of polysilicate loading upon the tensile mechanical response of 29 
day old THF-based PMAA-PTMO SIPNs. 









98 



80 
70 



* 60 



50 



CO 
CO 

CD 

c 

</5 40 

O) 

c 
a> 30 

CD 

c 
c 20 



LU 



10 - 



- 



/ 



fe: 



/ 



10 %/min strain rate 
L = 2.5cm 



Estimated Volume 
of Polvsilicate (%) 



1.6 

5.7 

12 

26 



_L 



_1_ 



50 






100 150 

% Elongation 



200 



250 



300 



Figure 4.9 The effect of polysilicate loading upon the tensile mechanical response of 52 
day old DMF-based PMAA-PTMO SIPNs. 



99 



Table 4.3 The effect of polysilicate loading upon the stress and elongation at failure of the 
PMAA-PTMO SIPNs for both co-solvent systems. 

Estimated Volume of Elastic Modulus Stress at Failure Elongation at Failure 

Polysilicate (%) (MPa) (MPa) (%} 

29 Day Old THF-Based SIPNs 

1.8 930,1170 34,35 410,360 

5.5 750±130 37±3.7 390±63 

11 870±71 39±1.5 300±24 
23 920±140 46±1.2 140±17 

52 Day Old DMF-Based SIPNs 

1.6 2150±140 51±1.1 26±5.5 

5.7 1070±98 41±3.0 270±30 

12 970±150 45±2.9 170±20 
26 980±280 40±3.5 27±10 



indicate that there is a definite thickness dependence in that exceptionally thick samples of 
both the DMF- and THF- based PMAA-PTMO SIPNs exhibit much higher yield strength 
and much lower elongation at failure than equivalent composition thin samples. For this 
reason only samples of reasonably similar thickness are displayed with the exception of the 
1.6% polysilicate DMF-based PMAA-PTMO SIPN, which illustrates this dependence. 

All of the inorganic loadings evaluated exhibit significant increases in the 
mechanical response as compared to the benchmark and y-irradiated gels. Indeed, greatly 
increased stress and elongation to failure along with yielding indicate significant changes in 
the morphology of the hybrids as a result of PMAA-PTMO SIPN formation Considering 
Table 4.3, there is no clear dependence of the elastic modulus upon the inorganic loading 
for either co-solvent system. However, for the THF-based gels there is the trend of 
increasing stress at failure and decreasing elongation at failure with increasing polysilicate 



100 

loading. Only the decreasing elongation at failure holds for the DMF-based gels when the 
thickness dependent response of the 1.6 vol.% polysilicate loaded SIPN is ignored. 

The stress and elongation at yield values are given in Table 4.4 along with the 
post-yield (P.Y.) stress drop and elongation prior to strain hardening, i.e., the percent 
elongation occurring after yield but before the onset of strain hardening. These post-yield 
values are depicted graphically in Figure 4. 10 as a function of polysilicate volume fraction. 
The data in Table 4.4 indicates that the stress at yield and elongation at which yielding 
occurs exhibit no correlation with inorganic loading. However, a clear trend of 
diminishing post-yield stress drop and post-yield elongation prior to strain hardening with 
increasing inorganic loading is observable in Figure 4. 10. Excluding the DMF-based 1 .6% 
polysilicate samples there is little difference between the post-yield stress drop of the two 
co-solvent systems. There is, however, a greater post-yield elongation prior to strain 
hardening for the THF-based gels. 



1 able 4.4 The ertec 


:t of polysilicate loading upon the tensile yield 


stress and elongation 


and post-yield (P.Y.) 


i response of the PMAA-PTMO SIPNs for both 


co-solvent systems. 


Estimated 


Stress at 


P.Y. Elongation 


Volume of 


Yield Elongation at P.Y. Stress Drop 


Prior to Strain 


Polysilicate (%) 


(MPa) Yield (%) (MPa) 


Hardening (%) 




29 Day Old THF-Based SIPNs 




1.8 


24, 26 3.7, 3.6 7.9, 8.6 


14, 10 


5.5 


19±1.4 4.3±0.90 4.7±0.48 


12±2.2 


11 


18±1.1 3.6±0.10 2.7±0.21 


6.7±0.25 


23 


22±2.5 4.0±0.22 None Observed 
52 Day Old DMF-Based SIPNs 


None Observed 


1.6 


76±3.1 5.4±0.12 25±2.1 


16*1.1 


5.7 


26±1.9 3.5±0.07 6.6±0.61 


17±1.7 


12 


24±3.5 3.7±0.14 2.2±0.84 


12±1.3 


26 


27±3.1 3.8±0.10 None Observed 


None Observed 



101 



* 30 
Q. 

Q. 
O 

5 20 

(A 

| 15 

CO 

s 10 

2 5 

v> 

o 

Q. 



L 



THF-basedPMAASIPNs 
(29 days old) 

DMF-based PMAA SIPNs 
(52 days old) 





5 10 15 20 25 

Estimated Volume of Polysilicate (%) 



30 



Figure 4.10 Post-yield stress drop and elongation prior to strain hardening as influenced 
by polysilicate loading and co-solvent system employed. 



102 

Traditionally, the phenomenon of yielding originates from the ability of a material 
to undergo viscous motion during applied stress and thereby dissipate energy. 148 This 
reorientation initiates by the formation of a neck in the sample. As the elongation 
increases, this neck grows at a rate dictated by the natural draw ratio of the material. This 
growth is generally accompanied by an increase in stress and modulus owing to the 
increased molecular orientation. Hence the term strain hardening is used to describe the 
process. In terms of the response of these SIPNs, the lower plot in Figure 4.10 indicates 
that the more spatially diffuse and polar THF-based polysilicate SIPNs are capable of 
undergoing greater elongation prior to the increase in stress attributable to complete 
reorientation. Considering the effect of the polysilicate loading, the observed decrease in 
both the post-yield stress drop and elongation prior to strain hardening for both co-solvent 
system indicates that as the continuity of the polysilicate phase increases, and presumably 
approaches a phase inversion, it increasingly hinders the ability of the SIPNs to undergo 
viscous motion. 
Dynamic Mechanical Response of the SIPNs 

Dynamic mechanical spectroscopy was performed on all the SIPNs of the same 
age with the storage modulus, E', and tan 6 response as a function of temperature 
appearing in Figures 4.1 1 and 4.12, respectively. Considering the storage modulus first, 
all 8 of the PMAA-PTMO SIPNs exhibit similar thermo-mechanical response. As the 
temperature increases from -145°C there is a gradual decay in the glassy modulus at ca. 
0°C for all samples. The Tg of PTMO gels containing no PMAA or polysilicate (other 
than the triethoxysilane crosslinks) is ca. -78°C. Therefore, the observance of this decay 
at much higher temperatures indicates good miscibility of the initial hybrid structure and 



103 



CD 
CL 

UJ 

o 



u 


a-ffff-i ' 


i 


1 

THF/IPA 


9 


SE&mmmEizz^ 




Estimated Volume 






8 


_ of Polysilicate (%): 

o 1.8 
a 5.5 

Q 11 






7 


v 23 




q&^ 



10 



CD 

a 

LU 
O) 

o 



9 - 



8 



o 1.6 


DMF/ 


* 5.7 




a 12 




v 26 








0.75°C/min 
- N 2 @ 200 ml/min 


i i^ 2 ^ 



-100 100 200 

Temperature (°C) 



300 



Figure 4. 1 1 Dynamic mechanical storage modulus, E', for both the THF- and DMF- 
based PMAA-PTMO SIPNs of the same age as a function of temperature for the 
polysilicate loadings indicated. 



104 



0.8 



0.6 



to 
g 0.4 



0.2 



0.0* 



N 2 @ 200 ml/min 
0.75°C/min 

Estimated Volume 
of Polysilicate (%): 

1.8 



THF/IPA 




0.8 



DMF/IPA 




-100 100 200 

Temperature (°C) 



300 



Figure 4.12 Dynamic mechanical tan 5 response for both the THF- and DMF-based 
PMAA-PTMO SEPNs of the same age as a function of temperature for the polysilicate 
loadings indicated. 



105 



the y polymerized PMAA. A plateau then follows till ca. 150°C at which point the most 
significant decay in the storage modulus begins. This decay is centered at ca. 200°C, 
which is close to the reported Tg of 230°C for PMAA. 149 Additionally, the magnitude of 
the decay is dependent upon the polysilicate loading and co-solvent system in that 
increasing levels of the inorganic phase reduce the decay associated with the Tg. Above 
the Tg of the PMAA phase the low volume percent polysilicate containing samples exhibit 
the thermally induced syneresis which has been observed in Chapter 3 and explained in a 
previous publication. 121 Interestingly, the effect of this syneresis is virtually undetectable 
in the samples containing more polysilicate. The exceptionally high value for the 0°C to 
150°C storage modulus region of the 1.8 vol.% polysilicate THF-based sample is 
indicative of the thickness effects discussed earlier. Conversely, the thinner 1.6 vol.% 
polysilicate DMF-based gel shown in the lower plot exhibits values consistent with other 
samples of the same relative thickness. 

Considering next the tan 8 response as a function of temperature displayed in 
Figure 4.12, multi-modal relaxations are observed in the spectrum of all samples 
investigated indicating that a wide variety of molecular environments exist within these 
gels. The overall breadth of the relaxations suggests that much mixing of the PTMO, 
polysilicate and PMAA phases is occurring. The intense relaxations centered at ca 1 90°C 
for all eight of the samples tested are attributed to the Tg of the PMAA phase. The 
increase in intensity of these relaxation with decreasing polysilicate loadings and hence 
increasing PMAA loadings suggest that the PMAA phase constitutes a significant portion 
of the molecular environment within these gels. This is further supported by the 
significant strength and modulus increases observed in the tensile samples. 



106 

Preliminary Investigations into SIPN Morphology 

Further evidence of the affinity of the PMAA for the polysilicate phase is provided 
by the atomic force microscopy images shown in Figures 4.13 and 4.14 of the fracture 
surfaces of both the TEOS(40) benchmark gel and TEOS(40) derived, 11 vol.% 
polysilicate loaded THF-based PMAA-PTMO SIPN, respectively. The tapping mode 
images clearly indicate that the tortuous, micro-phase separated polysilicate domains 
clearly observable in the benchmark gel are smoothed over in the SIPN gels. The 
morphological similarities of this polysilicate containing PMAA-PTMO- SIPN to those 
images collected by Toki et al. on the surface of poly(n-vinylpyrrolidone) (PVP)-silica 
hybrids are striking. 38 Their attribution of the homogeneity of the PVP-silica gel to 
hydrogen-bond formation between the PVP carbonyl groups and residual silanols, as 
measured using NMR, lends further support to the proposed poly(methacrylic acid)- 
polysilicate interactions discussed in this work. 

Small angle X-ray scattering was performed to evaluate the effect of SIPN 
formation upon the average electron density fluctuations present within the gels. Figure 
4.15 illustrates the changing scattering response observed upon formation of the SIPN 
within a benchmark 19 vol.% polysilicate containing gel. The scattering profile before 
swelling and y polymerization exhibits a maximum characteristic of these micro-phase 
separated gels. 30 An interdomain spacing of ca. 13 nm can be estimated from the 
reciprocal of the scattering vector at the maximum scattering intensity, I(s), for the 
starting gel. Upon polymerization of the adsorbed MAA the overall scattering intensity is 
reduced due to the lower volume fraction of the relatively electron rich inorganic phase. 



107 




X 0.200 UM/diu 
Z 50.000 nw/div 



Figure 4. 13 Atomic force microscopy image of the fracture surface of a 19 vol.% percent 
polysilicate, THF-based benchmark PTMO-silica gel collected using tapping mode. 



108 




X 0.200 PM/Jiv 
Z 40.000 rWdiv 



Figure 4.14 Atomic force microscopy image of the fracture surface of an 12 vol.% 
percent polysilicate, THF-based, PMAA-PTMO SIPN gel collected using tapping mode. 
This SIPN is formed from the same "parent" PTMO-silica gel as the piece imaged in 
Figure 4.13. 



109 



w 0.2 - 




0.08 



Figure 4.15 Small angle X-ray scattering profiles of the 19 vol.% polysilicate, THF-based 
benchmark gel and the subsequent 12 vol.% polysilicate, THF-based PMAA-PTMO SIPN 
resulting from y polymerization of the MAA swollen "parent" gel. 



110 

Additionally, there is no clear maximum in the scattering profile indicating that the PMAA 
and polysilicate phases have mixed and thereby reduced the electron density gradient. 
This data in conjunction with the DMS results presented above strongly suggests that 
mixing occurs in these PMAA-PTMO-silica SIPNs and is driven by the affinity of the 
PMAA and polysilicate phases for one another. 



Conclusions 





Significant changes in the tensile mechanical response of the benchmark PMTO- 
polysilicate hybrid composites can be induced by formation of PMAA-PTMO SIPNs. 
Additionally, the presumed spatial distribution and demonstrated polarity differences 
induced in the polysilicate phase by the use of two different co-solvent systems 
significantly affects the mechanical response of these y polymerized SIPNs. Overall, this 
transformation involves the conversion of elastomeric hybrids to high strength, high 
elongation organic glasses exhibiting yielding. The effect of increasing polysilicate loading 
upon the tensile mechanical properties of the SIPNs synthesized from both co-solvent 
systems has been found to increase the stress at failure, decrease the elongation to failure, 
decrease the magnitude of the post-yield stress drop and decrease the post-yield 
elongation prior to strain hardening. Within the individual co-solvent systems the THF- 
based SIPNs exhibit greater elongation than the near equivalently loaded DMF-based 
SIPNs. This indicates that less fully developed polysilicate networks are capable of 
undergoing greater viscous motion. It is unclear at this time whether the hydrogen 
bonding ability of the polysilicate phase directly influences the degree of viscous motion or 
is just an indicator of this less than fully developed glass network. Atomic force 



Ill 

microscopy and small angle X-ray scattering data suggest the formation of a PMAA rich 
SIPN phase adjacent to the polar polysilicate domains. Therefore, even if this hydrogen 
bonding does not directly contribute to the mechanical response, it does dictate the 
morphology of the SIPNs. 



CHAPTER 5 
THE EFFECT OF SOL CATALYST UPON OXYGEN DIFFUSION 

Relevant Background 

Motivation 

The potential of high temperature, gas separation processes in corrosive 
environments is driving the development of porous ceramic membranes. For these 
membranes to provide a separation mechanism, the pore diameters must be small enough 
for Knudsen diffusion to predominate. Although this mechanism is functional and has 
found use in the separation of uranium isotopes for the manufacturing of nuclear weapons, 
it is a slow process exhibiting relatively low gas selectivity. This selectivity can be greatly 
improved by using polymeric membranes. In general, rubbers exhibit the highest gas 
permeability of all the polymer classes. However, permeability and selectivity are virtually 
always inversely proportional 95 Therefore, although organic glasses such as 
polycarbonate, polyimide and polysulfone exhibit lower permeability than organic rubbers, 
the selectivity is improved. 113 ' 114 ,! 50 Surprisingly, little research has been performed 
on materials capable of true mixed-mode diffusion. Rather, the current interest is in 
understanding and modeling the dual-sorption characteristics of glassy polymers. 
However, an understanding of the gas transport mechanisms present in rubber/porous 
ceramic composites would provide a good foundation for further study of more refractory 
and corrosion resistant polymer-porous ceramic membranes combining the selectivity of 



112 



113 

polymers with the speed and additional selectivity of porous ceramics. 

Hybrid organic-inorganic composites offer a unique opportunity to study transport 
phenomena in composite systems. In Chapter 2, it was demonstrated that sol-gel-derived 
glass should exhibit Knudsen diffusion based on the pore diameters typically observed. 
Therefore, hybrids provide a single material in which two potential separation mechanisms 
are possible. The word potential is used because some controversy exists surrounding the 
presence of open and interconnected porosity within the inorganic domains of hybrid 
composites. Nevertheless, this unique morphological aspect paired with the other 
advantages that hybridization offers is stimulating interest in the utilization of hybrid 
composite technology for industrial gas separation processes, dissolved oxygen sensing 

devices, reduced pressure oxygen sensing paint and other niche applications. 79 > 8 ! > 90 ' ] 16 " 
120 

Although there is seemingly little interest in rubber-porous ceramic membranes, the 
PTMO-polysilicate system used in this study is free of many complexities present in the 
study of gas transport in organic glasses such as physical aging, free volume and quasi- 
ordered regions within the amorphous bulk. 89 The immediate goal of this work is to 
measure the effect of catalyst dependent morphology upon the oxygen diffusivity of these 
model hybrid composites. Furthermore, it is hoped that these results will provide an 
assessment of the presence of multiple diffusion mechanisms indicative of open, 
interconnected porosity within the inorganic polysilicate phase. 
Enabling Principle 

The measurement technique utilized is based on the detection of oxygen by a 
fluorescent compound embedded within the gels. The relationship between intensity and 



114 

oxygen concentration, as indicated by these molecular probes, is described by the Stern- 
Volmer relationship: 

-^ = l + k QM x M [0 2 ]. (5.1) 

1 [02] 

In this relation, I Q is the fluorescent intensity in the absence of oxygen, I[ 2] is the intensity 
at an oxygen concentration of [0 2 ], kq M is the bimolecular quenching constant and Tm is 
the lifetime of the luminescent compound in the matrix. 151,152 when the lumiphore 
dispersed within the sample comes into contact with molecular oxygen, the intensity at 
which it emits photons decreases. Therefore, this relationship predicts an increase in the 
intensity ratio from unity with increasing oxygen concentration. For materials such as 
liquids and rubbers, the concentration of a dissolved gas is directly proportional to 
pressure via Henry's law. The proportionality constant is the solubility, S. Therefore, the 
oxygen concentration, [0 2 ], can be replaced by the quantity SP02, where P02 is the partial 
pressure of oxygen. Since compressed air is used as the oxygen source for the 
experiments detailed in this chapter P O 2=0.21P and the solubility-pressure quantity can be 
substituted for [0 2 ] in Equation. 5. 1 to yield 

f = l + 0.21k QM T M SP. (5.2) 

p 






Therefore, the pressure-dependent phosphorescent intensity can be used to measure the 
concentration of molecular oxygen dissolved in the gel. Although a linear relationship is 
predicted, it often does not hold for polymers. The deviation from linearity is attributable 
to the influence of the micro-heterogeneous environment present in polymers upon the 
decay lifetime, x M , of the luminescent species. 82 For example, samples of lumiphore 



115 

coated silica particles dispersed within crosslinked polydimethylsiloxane as well as 
lumiphore doped polystyrene have been shown to exhibit two decay lifetimes each, albeit 
with different values. Although the resolved lifetime decays produce two linear Stern- 
Volmer responses in each of the samples over the pressure range of interest, the 
summation of the two linear responses leads to a nonlinear response for both samples. 82 
Therefore, the presence or absence of curvature in the Stern- Volmer plots can not be 
rigorously used as a check for the correct assumption of Henry's law. To ensure that the 
pressure-concentration response is indeed linear it would be necessary to perform detailed 
sorption experiments employing a micro-balance or pressure-volume apparatus such as 
that described by Koros et al.. 153 These experiments were not conducted in this study 
and, as stated above, Henry's law is assumed. As will be presented later, the oxygen 
diffusivity values obtained in this study are in good agreement with those published for 
other rubbers which suggests that the assumption is valid enough. 
Mass Transport Equation Utilized 

In the studies to be detailed shortly, lumiphore containing poly(tetramethylene 
oxide)-polysilicate hybrid composites were coated onto glass substrates. The rate of 
oxygen transport out of the exposed surface of the films is then detected by the change in 
oxygen concentration as assessed by the change in luminescent intensity. A special 
apparatus designed to do exactly this will be described in the Experimental section. The 
current section, however, provides the transport equations that will be applied to the data 
collected by this apparatus. 

The solution to any time-dependent diffusion process begins by solving Fick's 
second law: 



116 



ac_ p a 2 c 

at ax 2 



(5.3) 



for the time, t, dependent change in concentration, C, with position x. Under the 
assumptions of one-dimensional flow out of a homogeneous slab of infinite length and 
constant average diffiisivity, D, the separation of variables technique can be applied under 
the appropriate boundary conditions to produce the solution 



C-C 



Min 



C -C 

v< Max *■* Mm 



1-41 



7t 2 n=0 (2n + l) 2 



exp 



(2n + l) 2 Dt 
4 V 



(5.4) 



The derivation of this expression can be found in several texts on mass transport. 104 ' 154 
In this equation, L is the sample thickness. For the time-scales examined in these studies, 
the first iteration of the series is sufficient and produces the expression 



C-C 



Min 



C-Max C Min 



7t 



exp 



(5.5) 



where 



4L 2 



ti 2 D 



(5.6) 



and is referred to as the time constant. 154 The best fit of Equation 5.5 to a time- 
dependent change in concentration allows the calculation of x. Once x is known, the 
average diffusivity can be calculated. Assuming concentration is directly proportional to 
pressure within the film and taking into account the partial pressure of oxygen in air, then 
Cm;,, ■ 0.21SP M in, and C max ■ 0.21SP Ma x This allows Equation 5.5 to be rewritten as 



P - P 

r Max r Mii 



7t 



exp 



(5.7) 



117 

which is the form most conducive to these experiments since accurate pressure values are 

easily obtainable. 

Catalyst Effects on Polvsilicate Porosity 

Three different approaches have been taken to produce polysilicate phases 
exhibiting distinctly different continuity and levels of interaction with the surrounding 
polymer matrix. As the previous two chapters have elaborated upon, the choice of 
processing route and exposure of the gelling alkoxysilanes to different pH produces 
significant changes in the degree of organic/inorganic phase interaction. This chapter 
exploits these effects by first producing HC1 catalyzed gels which dynamic mechanical 
spectroscopy will reveal, possess the characteristic intimate level of mixing shown to exist 
in Chapter 3. This level of phase interaction is decreased somewhat by prehydrolyzing 
equivalent amounts of tetraethoxysilane in an HC1 acidified alcohol/water mixture for 2 
days prior to sol addition and composite gelation. Thus allowing the polysilicate 
"clusters" to grow in the absence of the polymer chains and to achieve a more spatially 
discrete form prior to consolidation of the composite. The limit in reduced phase 
interaction is achieved by using a base catalyst. The catalyst employed is the same 
ethylamine in water solution used in Chapter 3. In this last case, prehydrolysis will not be 
necessary since dynamic mechanical spectroscopy results will reveal that a significant 
reduction in phase mixing and interaction is observed in these gels. 

In addition to the phase interaction spectrum produced, it is anticipated that these 
processing variations will give rise to changes in the type of pore structures present. For 
example, Figure 5. 1 schematically illustrates the type of porosity that may be formed using 
the acid and base catalysis route. The HC1 catalyzed gels synthesized without 



118 




Acid catalysis 



Polymer 



Base catalysis 



• Polysilicate 



CK 



Figure 5.1 Schematic illustration of the types of porosity anticipated to be present as a 
result of the catalysts employed in this study. 



119 

prehydrolysis may possess polysilicate phases resembling that drawn in the upper left 
corner of the schematic. These pores, if present, are likely to be ill-defined, random 
defects within the polysilicate network resulting from incomplete hydrolysis and 
condensation. Consequently, they defy the traditional definition of a pore. Alternately, 
the ethylamine-catalyzed gels may possess inorganic phases resembling the drawing 
appearing in the upper right corner. The origin of the better defined pore structure would 
be the increased level of network development that is characteristic of base catalyzed 
systems and that drives enhanced phase separation. The increased level of phase 
separation may lead to larger agglomerates composed of smaller, more colloidal silica gel 
particulates. The packing defects associated with these clusters would give rise to a more 
traditional pore structure. By increasing the loading of polysilicate, the probability of 
reaching a porosity percolation limit is increased. Therefore, four polysilicate loadings 
have been investigated. 4.5, 11, 19 and 42% (vol/vol). Clearly stated, the objective of 
this study was to measure the diftusivity of the micro-phase separated composites as a 
function of both polysilicate loading and catalyst induced morphology. 

Experimental 

Triethoxysilane end fiinctionalized PTMO of molar mass 2,495 g/mole was used 
for all of the compositions produced in this study. The first batching that occurred was 
that of the prehydrolyzed TEOS and was accomplished by adding 0.19, 0.37 and 0.64 ml 
of TEOS by syringe to 3 ml polypropylene serum vials sealed with rubber stoppers. Each 
of the 3 vials contained a vigorously stirred isopropanol (IP A) solution that had been 
acidified with 0.87N HC1. The amount of IP A present was that necessary to produce a 



120 

volumetric ratio of 0.7:1 TEOS:IPA. Similarly, the amount of 0.87N HC1 present in the 
IPA prior to the addition of TEOS was the stoichiometric amount required to hydrolyze 
every ethoxy group of every TEOS molecule added. These ratios were held constant for 
each of the three sols produced. Each sol was then magnetically stirred for 2 days at 
ambient temperature and pressure. After 2 days, each of these sols were added to vials 
containing stirred, pre-batched solutions of end functionalized PTMO, IPA and the 
amount of a toluene and platinum (II) meso-tetra(pentafluorophenyl) porphine (PtTFPP) 
solution required to produce a lumiphore mass loading of 0.3% in the final gel. The 
luminescent compound PtTFPP was obtained from Porphyrin Products and used as 
received. The amount of TEOS added to each vial was that necessary to produce sols of 
20, 40 and 70% (mass/mass) with respect to the total mass of TEOS and PTMO. No 
additional catalyst was added. Each of these 3 sols were then magnetically stirred for 1 
minute prior to the deposition of a sol aliquot onto 1/2" x 5/8" numbered glass substrates 
that had been pre-weighed 3 times each. 

The processing of the remaining two sols was comparatively easy. Equivalent 
amounts of TEOS to those used in the prehyrolyzed sols were added via syringe directly 
to the PTMO, IPA PtTFPP/toluene solutions. After 5 minutes of stirring, the appropriate 
amounts of 0.87N HC1 or 70 wt.% ethylamine in water solution were added. The 
resulting mixtures were stirred for one minute then the appropriate sol aliquots were 
deposited onto the glass substrates. In the case of the HC1 catalyzed gels, 100% of the 
stoichiometric water needed for complete hydrolysis was added. However, for the 
ethylamine catalyzed gels only 47% of the required water was added. The deficiency in 
water is necessary to maintain optical transparency in the resulting gels. As was the case 



121 

for the prehydrolyzed TEOS samples, the amount of PtTFPP present in each sol was that 
necessary to produce a final mass loading of 0.3% within the gel. Samples produced 
without prehydrolyzing the TEOS are henceforth referred to as in-situ precipitated. Two 
additional gels were produced and function as baseline samples. These two samples were 
HC1 and ethylamine catalyzed gels that were each crosslinked by the sol-gel-processing of 
end functionalized PTMO without the addition of TEOS. 

To produce the sample films, the appropriate amount of sol was deposited onto the 
glass using a volumetric pipette after 1 minute of stirring. These glass substrates were 
double-sided taped to the bottom of 50 mm diameter polystyrene petri dishes. After 
deposition, the dish was covered and tilted to spread the sol over the entire glass surface. 
This was done quickly so that the dish could then be set on a level plate with sufficient 
time remaining for the sol to redistribute evenly over the surface. Each sample was then 
given 4 days covered in darkness to gel, although gelation typically occurred within hours. 
The absence of light was necessary to prevent photodegradation of the lumiphore. After 4 
days the covers were removed and residual solvent was allowed to evaporate from the 
gels. Next, the samples were placed into a 40°C vacuum oven and evacuated to 10 Torr 
under continuous pumping for 24 hours. The oven door was covered with aluminum foil 
to prevent the entry of ultraviolet light. Upon removal from the oven, the coated 
substrates were carefully detached from the petri dishes. Any adhesive remaining on the 
glass was removed using acetone and a tissue. All of the samples were then weighed 3 
times each to determine the mass of the coating. 

One final comment is due with regard to the volume of material deposited on the 
glass substrates. A spreadsheet has been developed which automatically performs the 



122 

calculation of the mass and volume of each component needed to produce sols with the 
desired catalyst equivalents, molar ratio of water-to-alkoxy and solids loading. This 
spreadsheet also calculates a theoretical density for gels of any composition, which is 
accomplished by scaling the density of the gelling species with the molar mass change of 
the condensing components. Agreement between predicted and measured densities is 
within 2% for gels containing 33% (vol/vol) polysilicate or lower. Knowing the density of 
the gel resulting from the sol deposited, it is possible to adjust the volume deposited on 
the substrate to produce a film of desired thickness once the surface area of the substrate 
to be coated is determined by mass, thickness and density measurements. Therefore, for 
this study, the volume of sol deposited has been adjusted to provide final coating thickness 
of 40 u.m for all polysilicate loadings. 

Two primary characterization techniques were employed for this investigation: 
dynamic mechanical spectroscopy (DMS) and a luminescence-based apparatus capable of 
measuring the pressure-dependent luminosity of oxygen sensitive films as a function of 
time. As with the DMS experiments described in Chapters 3 and 4, a heating rate of 
0.75°C/min was used from -150°C to 250°C. A nitrogen atmosphere was maintained in 
the furnace using a flow rate of approximately 200 ml/min. The frequencies tested were 
0.1, 0.5, 1, 5 and 10 Hz. All spectra displayed have been collected at 1 Hz. 

The apparatus used to measure the oxygen diffusivity of the films deposited on 
glass was designed and built in the Department of Aerospace Engineering, Mechanics and 
Engineering Science at the University of Florida and is under the supervision of Dr. Bruce 

Carroll.^ ' The apparatus, which is schematically illustrated in Figure 5.2, relies upon a 
photomultiplier tube to measure the pressure-dependent luminescence of polymer films 






123 



DC 



Photomultiplier Tube 




UVLamp 



Vacuum zzz^QQ 



Computer 




Mirror/Filter 



(£s) Compressed Air 



Sample Chamber 



Figure 5.2 Schematic illustration of the apparatus used to measure the time dependent 
intensity of luminescent gel samples. 



124 

coated on impermeable substrates. The samples are contained within a small sample 
chamber, and a 1.26 cm 2 diameter portion of the top surface of the film is exposed to gas. 
The atmosphere above the sample is maintained by two computer operated valves that 
control the pressure of a large vacuum reservoir and a compressed air tank. The chamber 
possess a single window parallel to the surface of the sample through which the incident 
ultraviolet light can pass along with the subsequent wavelength shifted luminescent 
emission. Not shown are the power supplies and amplifiers necessary for the apparatus to 
function. 

Although amplifier and power source settings have to be set manually, the 
software package Lab View® is used to control some operations of the apparatus such as 
the valve positions (open or closed). The software is critical for the collection of data, 
which is accomplished using a three-step procedure: 

1. A virtual instrument (VI) in Lab View® is run, and it records the actual 
pressure, as measured by a transducer, and the luminescent intensity of the 
sample, as measured by the photomultiplier tube (PMT) voltage, at 21 
different pressures between ca. 0. 1 and 1 . 1 atms. Thirty seconds is allowed 
for equilibration between each data point. Upon collection of the data, a 
nonlinear regression through the 21 points is automatically performed 
thereby calibrating PMT voltage to pressure above the sample and, 
consequently, oxygen concentration within the sample. 

2. A second VI is run that closes the "upstream" valve, thereby exposing the 
sample chamber to vacuum pressure of ca. 0. 1 atm. A period of 60 seconds 
then ensues to allow equilibration. After this minute has elapsed the VI 
simultaneously opens the upstream valve and closes the downstream valve 
thereby "instantaneously" increasing or stepping the sample to a higher 
pressure (typically 1.1 atm). A duration of 60 seconds passes during which 
the VI is recording the pressure in the sample chamber via the transducer 
and the PMT voltage, which is calibrated to concentration of molecular 
oxygen within the gel sample, as a function of time. For these experiments 
40 data points are collected each second for a total of 2,400 pressure/time 
data points. 



125 

3. The calibration VI is reopened and run again so that another calibration 
curve for equilibrium sample intensity (PMT voltage) as a function of 
pressure is obtained. The VI renames the "original" calibration data to 
backup files and then updates the current calibrations. 

The collection of two calibration curves allows a comparison of the equilibrium 
pressure/intensity response of the sample both before and after the step response. This is 
needed to ensure that a significant amount of photodegradation has not occurred which 
would invalidate the step response just collected. It should be noted that no significant 
change occurred in the pre- and post-step calibrations for the step responses presented in 
this work. 

The nonlinear equation regressed through the calibration data is 



1 

= a + bP + c 



dP T 

Trans 

. . =ati;rtt 

V PMT U + dP Trans 



r ah \ 

(5.8) 



where V PM t is the PMT voltage and P Tr an S is the transducer pressure. By allowing 30 
seconds for the samples to equilibrate prior to recording the PMT voltage, it is assumed 
that Pirans. is equal to the pressure of air within the gel sample. The constants a, b, c and d 
are determined during the regression. The general form of the equation is that of a linear 
pressure-voltage response with a Langmuir term added to account for any deviation from 
linearity that is sometimes observable in polymeric samples, as discussed earlier in this 
chapter. 

After the step response is collected, it is possible to curve fit a diffusion equation 
of the form given in Equation 5.7 using the graphical software package SigmaPlot*. This 
program uses a Levenberg-Marquardt least squares curve fitting routine to optimize the fit 
of the equation based on the parameters such as x. Additionally, the package also 
produces an uncertainty in the determination of each parameter. This uncertainty is 



126 

termed the standard error and is used in the calculation for oxygen diffusivity uncertainty 
discussed later in this chapter. 

Results and Discussion 

Film Thickness Measurements 

The thickness of all films used in this study is given in Table 5.1. As, the volume 
of sol deposited was intended to produce films 40 urn thick, in most cases, satisfactory 
agreement between the actual and target thickness was achieved. The error for all samples 
is one standard deviation of the mean arising from the three mass measurement technique. 
Although this standard deviation does account for the propagation of error involved, the 
visible rounding of the edges of the films deposited on the glass undoubtedly leads to an 
error exceeding that estimated from instrument uncertainty. Although more accurate 
techniques for film thickness measurement do exist, such as ellipsometry or spectrometric 
methods, the large sample surface area tested in the diffusion experiments, 1.26 cm , 
suggests that this average thickness technique may provide just as good of a value as these 

Table 5. 1 Estimated thickness for the samples used in oxygen diffusivity measurements. 

Sample Polysilicate Loading 

HC1, 100% H 2 0, in-situ TEOS 4.5 
HC1, 100% H 2 0, in-situ TEOS 1 1 
HC1, 1 00% H 2 0, in-situ TEOS 1 9 
HC1, 100% H 2 0, in-situ TEOS 42 
HC1, 1 00% H 2 0, prehy. TEOS 1 1 
HC1, 100% H 2 0, prehy. TEOS 19 
HC1, 100% H 2 0, prehy. TEOS 42 
Ethyl, 47% H 2 0, in-situ TEOS 4.5 
Ethyl., 47% H 2 0, in-situ TEOS 1 1 
Ethyl., 47% H 2 0, in-situ TEOS 1 9 
Ethyl., 47% H 2 Q, in-situ TEOS 42 



Thickness (um) 


39±0.2 


43±0.3 


43±0.2 


50±0.2 


46±0.3 


43±0.2 


47±0.2 


44±0.2 


46±0.3 


46±0.2 


53±0.2 



127 

point source measurements. For the estimations of error used in this chapter, an 
uncertainty of 5% of the film thickness will be used. 
Dynamic Mechanical Spectroscopy 

DMS was performed on gels containing 33 vol % polysilicate to confirm the 
difference in the phase interaction anticipated to be present in gels produced using the 
above detailed sol processing. The level of phase interaction is indicative of the spatial 
distribution of the polysilicate phase. Figure 5.3 displays the storage modulus, E', (top 
portion) and loss dispersion, tan 5, response (bottom portion) of each gel. Considering E' 
first, all gels exhibit nearly identical response until ca. -75°C which is the glass transition 
temperature of the PTMO phase. The decay in modulus occurs at approximately the same 
rate for all three samples until 0°C where significant changes are observable. As discussed 
in Chapter 3, much information can be gained regarding the level of phase interaction 
occurring from the rubbery regime of these hybrid composites. The fact that the HC1- 
100% H 2 gel possesses the highest rubbery modulus indicates that it also possesses the 
highest degree of phase interaction as entanglements with the polysilicate phase behave as 
labile crosslinks. An increase in modulus occurs beyond 175°C indicating the lack of a 

completely developed polysilicate network. 12 1 The lowest rubbery modulus is observed 
for the ethylamine catalyzed gel utilizing 47% water. The reduced modulus indicates a 
reduced level of phase interaction, and the absence of any thermally induced modulus 
increase indicates the development of a nearly complete polysilicate network despite the 
water deficient sol the gels were produced from. The prehydrolyzed, HC1 catalyzed 
TEOS sample exhibits a rubbery response containing features of both the HC1 and 
ethylamine catalyzed gels. At 25°C, the gel possesses a modulus nearly identical to that of 



128 



? 8.5 



7.5 L 
0.20 

0.15 f 



0.00 



T 1 1 

N 2 @ 200 ml/min 
0.75 °C/min 




HCI-100%H 2 O 

HCI Pre-TEOS 
100% H 2 

Ethyl.-47%H 2 0_ 




-150 -100 -50 50 100 150 200 250 

Temperature (°C) 



Figure 5.3 Dynamic mechanical storage modulus and tan 5 response of gels containing 33 
vol.% polysilicate derived from sols employing acid and base catalysts 



129 

the HC1 catalyzed, in-situ precipitated sample. As the temperature increases, the modulus 
continues to decay to a minimum very close to that of the ethylamine catalyzed gel. 
However, the syneresis typical of acid catalyzed gels is observable. 

Considering the tan 8 response, significant damping begins to occur at the onset of 
the Tg at ca. -75°C and continues until ca. 100°C for all three gels. The overall breadth of 
the relaxation spectra indicates that a wide variety of molecular environments exist within 
these gels. The most uniform or normally distributed relaxation occurs for the HC1 
catalyzed, in-situ precipitated polysilicate. This indicates that a high degree of phase 
mixing is occurring. Conversely, a bimodal distribution is observed for the ethylamine- 
catalyzed gel with peaks centered at ca. -50°C and 10°C. This bimodality indicates the 
presence of two phases with each being predominantly richer in one species than the other. 
The acid catalyzed, prehydrolyzed gel exhibits a fairly uniform distribution. However, a 
low temperature shoulder does exist at ca. -45°C indicating the presence of some phase 
segregation. Again, this suggests that an intermediate level of mixing is achieved. With 
regard to damping, i.e., the maximum value of tan 5, the prehydrolyzed gel exhibits the 
highest propensity to absorb energy. It is unclear at this time why the intensity does not 
also fall in an intermediate position. However, the above analysis of the E' response as 
well as the breadth and shape of the tan 6 responses indicate that three varying levels of 
phase separation have been achieved. 
Oxygen Diffusivity 

The step responses of the HC1 catalyzed, in-situ precipitated gels are shown in Figure 5.4, 
while the responses of the HC1 catalyzed, prehydrolyzed TEOS and ethylamine catalyzed 
gels are displayed in Figures 5.5 and 5.6, respectively. Note that the ordinate of 



130 



■M 1.2 



X 

03 




Volume of 
Polvsilicate (%) 



4.5 
11 
19 
42 



20 30 40 

Time (sec) 



50 



60 



Figure 5.4 Step responses of the four HCl catalyzed gels utilizing the in-situ precipitation 
of polysilicate in the presence of 100% of the stoichiometric water required for hydrolysis. 









131 




Volume of 
Polvsilicate (%) 



4.5 
11 
19 
42 



20 30 

Time (sec) 



40 



50 



60 



Figure 5.5 Step responses of the four HCl catalyzed gels utilizing prehydrolysis of TEOS 
in the presence of 100% of the stoichiometric water required for hydrolysis to produce 
polysilicate clusters prior to sol batching. 



132 




10 



Volume of 
Polvsilicate (%) 



4.5 
11 
19 
42 



20 30 40 

Time (sec) 



50 



60 



Figure 5.6 Step responses of the four ethylamine catalyzed gels utilizing in-situ 
precipitation of polysilicate in the presence of 47% of the stoichiometric water required 
for complete hydrolysis. 



133 

each of the samples is vertically shifted by 0.2 to make the response of each clearly visible. 
Only at the maximum loading of 42% (vol/vol) is any visually significant change in the 
response observed for the either class of acid catalyzed gels (Figures 5.5 and 5.6). 
Similarly, it can be seen that the response of the ethylamine catalyzed gels is different than 
that of the acid catalyzed gels. However, within this class, all four loadings of the 
ethylamine catalyzed gels investigated exhibit near identical responses. 

As the thickness of all the samples is nearly equal, it is possible to examine the 
influence of processing by regrouping the step responses according to their polysilicate 
loadings. This has been done, and the groupings appear in Figures 5.7 through 5.10. The 
polysilicate loadings for each graph are given in the upper right hand corner. Also shown 
in these figures are the modeled responses to be discussed shortly. 

Disregarding the modeled fits momentarily, a review of Figures 5.7 through 5.9 
indicates that the in-situ precipitated, HCL catalyzed samples respond most quickly to the 
pressure change for all loading up to 19 vol % polysilicate. Conversely, the response is 
slowest for the in-situ precipitated, ethylamine catalyzed samples. The prehydrolyzed gels 
exhibit an intermediate time response. Beyond 19% (Figure 5.10), a reversal is observed 
in that the ethylamine catalyzed gels now exhibit the fastest response. 

The exact model fit to the step responses shown above is 



° °Min 

°Max ~ Pfvlin 



= B - C exp 



(5.9) 



The expected value of B should be unity, and the expected value of C should be S/n 2 
(0.81). The parameter values, associated uncertainties and calculated oxygen diffusivity, 
D 02 , for each sample are given in Table 5.2. Visual examination of the curve fits in 



134 



1.2 

1.0 

cl 0.8 

X 

■ 

0- 0.6 



Q_ 

■ 0.4 

Q. 



0.2 
0.0 



10 



20 



4.5 Vol. % Polysilicate 




HCI-100%H 2 O 
Ethyl.-47% H 2 
Diffusion Model 



30 
Time (sec) 



40 



50 



60 



Figure 5.7 Step responses and diffusion model curve fits for HCL and ethylamine 
catalyzed gels containing PtTFPP and an estimated 4.5 vol.% polysilicate. 






135 



1.2 
1.0 

c 

o? 0.8 \ 
■ 

^ 0.6 



.£ 
Q. 



0.4 \ 

0.2 

0.0 



I I I 1 1 1 

1 1 Vol. % Polysilicate 


§ f o HCI-100% H 2 

f / Q HCI Pre-TEOS 
M 100% H 2 

M a Ethyl.-47% H 2 

ff Diffusion Model 



10 



20 30 40 

Time (sec) 



50 



60 



Figure 5.8 Step responses and diffusion model curve fits for the in-situ precipitated and 
prehydrolyzed HCL catalyzed gels, as well as the ethylamine catalyzed gels. All samples 
contain PtTFPP and an estimated 1 1 vol.% polysilicate. 



136 





1.2 




1.0 


c 

is 

Q. 


0.8 


X 

CO 

5 
0_ 


0.6 


£ 
5 
0_ 
i 


0.4 


0_ 






0.2 




0.0 



10 



19 Vol. % Polysilicate 




HCI-100%H 2 O 

HCI Pre-TEOS 
100% H 2 

Ethyl. -47% H 2 

Diffusion Model 



20 30 

Time (sec) 



40 



50 



60 



Figure 5.9 Step responses and diffusion model curve fits for the in-situ precipitated and 
prehydrolyzed HCL catalyzed gels, as well as the ethylamine catalyzed gels. All samples 
contain PtTFPP and an estimated 19 vol.% polysilicate. 



137 



1.2 



1.0 - 



of 0.8 - 



°- 0.6 

C 

5 
EL 



■ 0.4 - 



0.2 
0.0 



i i i i r - ■ 1 

42 Vol. % Polysilicate 


Jf ^s$*^ ° 


HCI-100%H 2 O 


# jtdSr^ a 


HCI Pre-TEOS 
100% H 2 


flf A 


Ethyl.-47% H 2 


__g i i i 


- Diffusion Model 



10 



20 30 

Time (sec) 



40 



50 



60 



Figure 5.10 Step responses and diffusion model curve fits for the in-situ precipitated and 
prehydrolyzed HCL catalyzed gels, as well as the ethylamine catalyzed gels. All samples 
contain PtTFPP and an estimated 42 vol.% polysilicate. 



138 

Figures 5.7 through 5.10 indicates good agreement with the exception of the 42% HC1 
catalyzed gels. These two samples exhibit step responses during the first 3 seconds that 
can not be described by this model. The most likely explanation is that insufficient time 
has been allowed for the samples to reach equilibrium luminescence. This is supported by 
the estimated B values of 1 .27 and 1 .28. 

Although the parameter uncertainties are automatically estimated by the curve 
fitting routine, the uncertainty associated with the diffusivity measurements in Table 5.2 
must be estimated using an uncertainty analysis of the relevant parameters. Equation 5.6 
can be written in terms of the diffusivity to yield: 
4L 2 

The uncertainty in this measurement can be estimated from the summation of the partial 
derivatives with respect to thickness and x 



dD 02 



8L 

7TT 



^4L 2 



w. 



V7CT 2 



(5.11) 



Table 5.2 Diffusion model parameters and resulting oxygen diffusivity. 



Sample 


B 


C 


x (sec) 
2.13±0.015 


D 02 ( 10 6 cm2/sec) 
2.9±0.30 


4.5%, HC1, in-situ 


0.991 


0.874±4.18E-3 


11%,HC1, in-situ 


1.01 


0.982±2.64E-3 


2.63±0.011 


2.8±0.28 


19%, HC1, in-situ 


1.00 


0.972±1.32E-3 


5.81±0.013 


1.3±0.13 


42%, HC1, in-situ 


1.27 


1.13±3.84E-3 


43.4±0.31 


0.24±0.024 


ll%,HCl,prehy. 


0.995 


0.979±2.01E-3 


3.32±0.010 


2.6±0.26 


19%,HCl,prehy. 


0.996 


0.986±1.20E-3 


6.50±0.013 


1.1*0.11 


42%, HC1, prehy. 


1.28 


1.15±3.91E-3 


43.6±0.31 


0.20±0.021 


4.5%, Ethyl., in-situ 


0.993 


0.961±1.47E-3 


7.33±0.019 


1.1*0.11 


11%, Ethyl, in-situ 


0.996 


0.982±1.37E-3 


6.94±0.017 


1.3*0.13 


19%, Ethyl., in-situ 


0.997 


0.980±1.03E-3 


9.09±0.018 


0.92±0.093 


42%, Ethyl., in-situ 


0.994 


0.967±1.12E-3 


10.1±0.023 


1.1*0.11 



139 

where L is the sample thickness, and w L and w T are the uncertainties in the values of the 
thickness and x, respectively. 155 ' 156 Although a measure of error does exist for the 
thickness of each gel and is given in Table 5.1, a ballpark uncertainty of 5% is assumed to 
allow for the assumptions made in the film thickness calculation. Since the uncertainty of 
x for each gel has been given in Table 5.2, it is possible to calculate the oxygen diffusivity 
uncertainty for each sample. An assessment of the resulting uncertainty for these D02 
values indicates that ca. 10% error exists in the measurements. 

The graphical display of the oxygen diffusivity as a function of polysilicate loading 

is given in Figure 5.11. These results reveal that the oxygen diffusivity of the acid 

catalyzed, unfilled gel, i.e., crosslinked in the absence of additional TEOS, is ca. 3x1c 6 

cm 2 /sec while that of the base catalyzed, unfilled gel is three times lower. The first check 

that must be made is the validity of these values. Since no published values of the oxygen 

diffusivity of PTMO can be found, comparisons must be made to other rubbers. 

Poly(dimethyl siloxane) (PDMS), which is the most permeable rubber known, possesses 

an oxygen diffusivity estimated by a similar technique at 3.55xlO* 5 cm 2 /sec at 25°C. 152 

Additionally, values of 1.75XK)- 6 and lSxlO -6 cm 2 /sec at 25°C for natural rubber and 

polybutadiene, respectively, have been published. 157 That good agreement is found 

between the oxygen diffusivity of these unfilled composites (both acid and base catalyzed) 

and these published values of natural rubber and polybutadiene is encouraging. This 

validity is enhanced by the fact that all of these published values are less than that of 

PDMS measured using a similar fluorescence technique. ] 52 Having established that the 

values measured are reasonable, it becomes necessary to explain why the acid catalyzed 



140 



o 

CD 

E 
o 

"o 



CM 

o 



> 

i 

b 

c 

0) 



T 



"T 



O- HCI-100%H 2 O 

D-- HCI Pre-TEOS 
100% H 2 

a- Ethyl.-47% H 2 




4 



14 21 28 35 

Volume of Polysilicate (%) 



42 



Figure 5.11 Oxygen diffusivity as a function of polysilicate loading for all three classes of 
gels produced. 



141 



unfilled gels, i.e., 4.5 vol.% polysilicate, has a diffusivity that is different from the base 
catalyzed gel. No complimentary data exists for these 4.5 vol.% polysilicate gels. 
However, insights can be gained from examination of more highly loaded gels, e.g. 33 
vol.%. A review of Figure 5.3 reveals that the storage modulus at 25°C for the ethylamine 
catalyzed gel is lower than that of the HC1 catalyzed gel. This reduced value is attributed 
to the diminished reinforcing effect induced by the more phase separated, base catalyzed 
polysilicate phase. This conclusion is reinforced by the fact that swelling of both the HC1 
and ethylamine catalyzed, in-situ precipitated gels in tetrahydrofuran for 24 hr results in 
mass uptakes of 19.8±1.22% and 31.1±2.60%, respectively. Therefore, the ethylamine 
catalyzed gel likely possesses a lower crosslink density which should lead to increased 
diffusivity. Figure 5.11 indicates that this is indeed observed. It is reasonable to assume 
that this trend of decreased phase interaction leading to increased diffusivity would hold 
for all polysilicate loadings generated using base catalysis Consequently, the 4.5 vol % 
ethylamine catalyzed gel should still possess a greater diffusivity than the HC1 catalyzed 
gel. However, the opposite is observed. This indicates that the slight variation in 
measured diffusivity may be due to the environment the lumiphore resides. Published 
results indicate that this is quite possible. 82 - 158 ' 159 The result of a slight change in the 
pressure-luminescence response would be slightly different step responses. The values of 
x calculated would then be different, as would the calculated diffusivity. Recall that 
exposure of HC1 catalyzed gels to ethylamine in Chapter 2 resulted in a significant 
reduction in the silanol concentration within the gels. It is quite possible that the in-situ 
precipitation of the polysilicate phase in the presence of ethylamine would result in a 
similar reduction in the number of silanol species present in the inorganic phase. 






142 

Therefore, the difference in measured diffusivities may reflect this change in chemical 
environment within the two gel classes. However, within each class the influence of 
polysilicate environment remains constant and the luminescent characteristics of the probe 
should remain the same. 

Considering next the influence of processing, a significant difference is observable 
between the acid and base catalyzed gel classes as the polysilicate loading increases. The 
diffusivity of both classes of the HC1 catalyzed gels decreases with increased polysilicate 
loading while the ethylamine catalyzed gels exhibit virtually no change. Few studies have 
been published detailing transport phenomena in hybrids because the emphasis has been on 
device production and performance. Insights can be gained, however, by adapting the 
analysis used to explain both the influence of crystallinity, phase separation and 
impermeable particles, such as zeolites, upon diffusion characteristics. 106,108,109,160 
Although different symbols and accompanying terminology are used to express the effect, 
the overall result is that the diffusivity of crystalline polyethylene, SBS block copolymers, 
filled rubbers and zeolite embedded polydimethylsiloxane decreases by a factor of k 
according to the relation 

__ amorphous 

D = " (5.12) 

where D amorphous is the diffusivity of the amorphous or "high diffusivity" region. The term k 
is essentially interpreted as the geometric impedance factor intended to account for the 
tortuosity of the diffusion pathways resulting from the presence of the less permeable or 
absorptive domains. In the polyethylene studies, it was demonstrated that at 22 vol.% 
crystallinity, the oxygen diffusivity is 1/10 of that estimated for amorphous 



143 



polyethylene. 160 In the case of polydimethylsiloxane containing dispersed zeolites, it was 
demonstrated that the presence of these gas immobilizing fillers delays the time necessary 
to reach the steady-state permeation required by the time-lag technique. However, only a 
1/4 reduction in the oxygen diffusivity of the composite system occurred at a 21 vol.% 
loading. 106 This is a significant difference in the tortuosity, or K-value, for these two 
types of fillers and indicates that as the oxygen absorptive ability of the second phase 
increases the impact it has upon reduced diffusivity is diminished. The minimum pore 
diameter present in the embedded zeolites was 0.45 nm. It is anticipated that the pores 
present within the polysilicate domains of these hybrids could be as much as an order of 
magnitude larger. Therefore, it is not unreasonable to conclude that the effect of such an 
absorptive second phase may be insignificant. As Figure 5.13 reveals, the ethylamine 
catalyzed gels do not exhibit any change in oxygen diffusivity with increased loading 
suggesting that the polysilicate domains present do contain significant porosity It is 
possible that the diffusivity of these pores is on the order of the polymer matrix they reside 
in. For example, the published value of the diffusivity of methane at ambient temperature 
in a single zeolite possessing 0.56 nm minimum pore diameter is 8x1c 6 cm 2 /sec. 100 This 
value is quite close to the measured lxlO" 6 cm 2 /sec for the hybrids. Therefore, if the pores 
present in the base catalyzed polysilicate phases are of comparable size then increased 
polysilicate loading would have no effect. In relation to Equation 5.12, this would 
constitute k equal to unity. 

The HC1 catalyzed gels exhibit a much different response and undergo significant 
reductions in the composite diffusivity with increasing polysilicate loadings. This suggests 



144 

that these domains contain less absorptive ability, such as smaller pores, and are much 
more impenetrable to the diffusing oxygen. The k values for increasing polysilicate 
loadings are given in Table 5.3 and reflect the increasing tortuosity of the diffusion 
pathways induced by the comparatively impenetrable polysilicate phase. Interestingly, 
there is very little difference in diffusivity induced by prehydrolyzing the TEOS compared 
to in-situ precipitation. This is most likely due to the similar "effective crosslink density" 
indicated by the near equivalent DMS storage modulus values at 25°C in Figure 5.3. 

One method of assessing the porosity present in the polysilicate domains of the 
gels is to force-fit a double-exponential diffusion equation of the form 

P-P ( t > i 

K ^- = B- C.exp -— +C,exp 

v V v x 2 y 

to determine the diffusivity of both the PTMO and polysilicate phase. Despite the fact that 
very good agreement already exists for the single-exponential model, this was done for the 
19 vol.% ethylamine catalyzed gel. Only one gel was needed as the diffusivity is 
independent of the loading. The resulting fit is not shown as no discernible difference in 
the quality of the single-exponential and double-exponential fits is observable. The 
resulting D 02 values are 1.2±0.24xl0" 6 cm 2 /sec and 0.63±0.18xl0^ cm 2 /sec for the PTMO 
and polysilicate phases, respectively. The diffusivity of oxygen undergoing transport in 



Table 5.3 Tortuosity factors as a function of increasing polysilicate loading for the HC1 
catalyzed gels. 



P -P 

1 Max ^Min 



P 



(5.13) 



Acid Catalyzed Polysilicate Volume (%) 



11 1.0 

19 2.2 

42 12 



145 

the Knudsen regime is given as 

i 
4rf2R.lV 

DK= TlisrJ' < 514 > 

For an assumed D K of 1.2x10"* cm 2 /sec the calculated pore diameter is ca. 4 x 10" 4 nm. 
This value is nonsensical, and suggests that although diffusion is occurring within the 
polysilicate phase it is most likely occurring in transport pathways that are considerably 
smaller than those necessary for Knudsen diffusion. Since the size of the pores strongly 
influences diffusion, these hybrids must possess pore diameters that are very near the 
collision diameter of the gas, i.e., 0.2 nm to 1 nm. Furthermore, the acid catalyzed 
domains are smaller and less well defined. The review of the literature concerning hybrid 
membranes given in Chapter 2 strongly supports these observations in that near universal 
absence of Knudsen diffusion was observed. 1 16-1 19 

Conclusions 

A luminescence-based technique for measuring the oxygen diffusivity of 
homogeneous materials has been applied to hybrid organic-inorganic composites. The 
near molecular level of mixing present in these sol-gel-derived composites has been tuned 
to produce three differing levels of polysilicate structure through the use of both acidic 
and basic catalysts as well as the prehydrolysis of the silicon alkoxide during sol 
processing. The results of this testing indicate that the technique is valid since the oxygen 
diffusivity values calculated from the application of a single-exponential diffusion model 
appear to be in agreement with the published results on other types of rubbers. Some 



146 



indication is present, however, that the molecular environment that the lumiphore probes 
reside within influences the measured oxygen diffusivity. 

One of the goals of this work was to assess the potential for well-defined, open 
porosity within these hybrid gels. It was anticipated that if porosity was indeed present 
the diffusion would likely be occurring in the Knudsen regime. The results collected in 
this study indicate that this is not the case. Although the polysilicate phases do possess 
some absorptive nature, with the ethylamine catalyzed gels exhibiting the highest 
propensity to absorb and/or immobilize oxygen, theoretical Knudsen diffusion models are 
unable to accurately predict realistic pore diameters. Therefore, the pore structures 
present in these hybrid composites most likely exhibit a configurational influence upon the 
diffusing oxygen, i.e., the types of pores present have dimensions approaching the collision 
diameter of the diffusing oxygen. 









CHAPTER 6 
CLOSING REMARKS 

Rubber Elasticity and Nonideal Networks 

In Chapter 3, the average molar mass between crosslinks was estimated using 
elementary rubber elasticity theory in conjunction with DMS. The resulting values were 
shown to be in excellent agreement with those obtained using traditional swelling studies. 
These results reveal that under the conditions of low strain amplitude and in the absence of 
thermally induced chemistry/curing, the combination of dynamic mechanical data and 
rubber elasticity theory are complimentary and do provide seemingly accurate values of 
network parameters. Perhaps, the most important word is complimentary. DMS is a 
widely used tool for evaluating the modulus, Tg, damping characteristics and phase 
homogeneity of polymers. Frequently, thermo-mechanical scans are performed on 
thermosets, and these parameters are tabulated. Less frequently, however, is additional 
information extracted from these spectra. The investigation detailed in Chapter 3 reveals 
that application of an idealized theory to a nonideal system can provide much additional 
information concerning phase mixing. Although the results of the analysis appear to be 
valid for this system, some would surely question the credibility of the approach. 
However, other evidence attesting to the success of rubber elasticity theory in evaluating 
nonideal networks exists. 

147 



148 

Although the elasticity of bimodal networks has been analyzed extensively over the 

past decade, these investigations have exclusively utilized the Mooney-Rivlin equation J"' 
Only recently have preliminary findings of swelling-based elasticity studies been reported. 
However, like elastomeric hybrid composites, the results suggest that these nonideal 
networks behave in a much more ideal way than expected. In many respects, the two 
systems are similar. For example, bimodal networks are comprised of exceptionally short 
oligomers and higher molar mass chains, e.g., 500 and 10,000 g/mole, respectively. 
Current thinking on these networks is that the short chains coalesce and form macro-sized 
crosslink junctions spanned by the longer chains. By analogy, the hybrids discussesd in 
this work possess macro-sized domains rich in polysilicate, a highly crosslinked network. 
These domains are the crosslink junctions spanned by the PTMO. When these results are 
taken as a whole, they suggest that rubber elasticity theory is quite capable of evaluating 
these heterogeneous networks, and it provides a very useful tool for examining phase 
mixing. 

However, several important issues remain to be answered. One of the underlying 
assumptions in the analysis used in Chapter 3 concerns the use of dynamic moduli data in 
place of equilibrium modulus values. The assumption made in applying the dynamic data 
is that if the characteristic relaxation time of the material is less than the time-scale of the 
DMS measurement then essentially the material is at equilibrium. The time required to 
obtain the dynamic modulus is approximated by the frequency of the applied strain. At 0. 1 
Hz, the time required for testing is 10 seconds. Several attempts have been made to 
determine the characteristic relaxation times for these gels. Although not reported on in 
Chapter 3, the results are interesting in that virtually no relaxation is observable at room 



149 

temperature. However, upon heating to 70°C, significant stress relaxation/retardation 
occurs. Attempts at using these equilibrium modulus/temperature data points to calculate 
the average number of covalently bonded, elastically active chains proved unsuccessful 
owing to the low number of data points and large scatter in the values. A better designed 
and more rigorous analysis of the stress relaxation and temperature response of these gels 
as a function of both polysilicate loading and spatial distribution would provide valuable 
insight into the extent of phase mixing occurring This analysis could then be compared to 
the network parameters measured using the DMS-based technique and swelling-based 
assessment of the same parameters. Taken as a whole, this would constitute an excellent 
study of the network homogeneity and ideallity present in hybrid composites. No such 
studies have been published in the hybrid community. 

High Performance SIPNs 

The origin of the investigation detailed in Chapter 4 was the high strength, high 
elongation and yielding previously observed for SIPNs containing poly(methacrylic acid) 

(PMAA) as the "thermoplastic" interpenetrating phase. ■ 23 The classification of these 
materials as semi-interpenetrating polymer networks is based on the fact that the PTMO 
and polysilicate hybrids from which the "SIPNs" are produced are pre-established 
thermoset networks. Swelling this network with an acrylic monomer and then 
polymerizing it in-situ produces the thermoplastic interpenetrating chains. Realistically, 
the high energy of y radiation induces chain scission and leads to increased modulus and 
strength and decreased elongation at break in the elastomeric hybrid gels. Therefore, 
when the monomer swollen gels are exposed to the radiation, it is quite probable that 



150 

some methacrylic acid segments or PMAA oligomers are incorporated into the PTMO 
chains. Furthermore, some crosslinking may be induced in the PMAA. These possibilities 
were part of the reason for addressing the issue of polysilicate loading and spatial 
distribution in the SIPNs. By varying the amount and distribution of this reinforcing 
phase, insights into the factors contributing to the high strength and elongation were 
identified. 

The results of Chapter 4 revealed that as the loading of polysilicate increased the 
elongation at failure decreased without a significant gain in strength. Similarly, when the 
degree of hydrogen bonding likely occurring was reduced by replacing THF with DMF in 
the sol, the elongation at failure decreased without a significant gain in strength. 
Accompanying the decreased elongation values were decreased post-yield elongation prior 
to strain hardening. These trends indicate that as the polysilicate loading and connectivity 
increases the glassy network is capable of less deformation under the load Therefore, 
since the use of an acid catalyzed or THF-based sol promotes the development of a less 
connected polysilicate phase, it is the ability of this network to deform that allows the high 
elongation. Inherent in this finding is that the unhindered PTMO-PMAA phases are 
capable of significant deformation. 

Several additional studies would compliment these findings well. For example, 
obtaining the mechanical tensile response of PMAA would provide a baseline against 
which the response of the hybrid SIPNs could be compared. However, efforts to 
compression mold PMAA of ca. 100,000 g/mole polymer have produced brittle glasses 
from which tensile samples could not be obtained. It is arguable that this relatively low 
molar mass is below the critical molar mass for entanglement and that this explains the 



151 

brittle nature. This is quite possible. Until higher molar mass polymer can be produced or 
synthesized, this baseline response can not be obtained. One advantage that this lower 
molar mass material may provide is enhanced diffusion rates. An interesting study would 
be to dissolve the 100 kg/mole PMAA in an alcohol and then swell pieces of the hybrid in 
this solution. If a significant amount of PMAA can be absorbed, it would result in a more 
ideal SIPN. Therefore, it would be possible to investigate the influence of the organic 
phase connectivity, or lack of, upon the mechanical response of the hybrids. Such a study 
would compliment the investigation of inorganic phase connectivity upon mechanical 
response. 

Gas Transport in Hybrid Composites 

The origin of this study was the distinct lack of information addressing gaseous 
diffusion in hybrid composites composed of a significant amount of the organic phase. 
This is in contrast to the organically modified glasses described in Chapter 2. The central 
theme of the study was to measure the diffusivity of a common, inert gas as a function of 
the polysilicate loading and structure. Again, it is the wide-variety of morphologies that 
can be produced in hybrid composites that make them such ideal models for many studies 
of mixed phase systems. In this study, it was hoped that the diffusivity values measured 
and trends observed would provide an assessment of the types of pore structures present 
in the polysilicate phase. The ability to detect significant porosity would substantiate the 
notion that pores do exist in these composites. 

The results were highly informative. Initially, it was thought that if pores do exist 
in the polysilicate phase, then Knudsen diffusion would predominate, as is the case in sol- 






152 

gel-derived membranes. Therefore, hybrid composites may possess two gas separation 
mechanisms. Evidence of this mixed-mode diffusion would have been step responses that 
could only be fit by a dual-exponential diffusion model. However, in agreement with the 
majority of literature published on organically doped sol-gel membranes, Knudsen 
diffusion was not observed. Rather, the materials exhibited diffusion responses indicative 
of configurational transport at best. Valuable insights were gained, however, from the 
trends in diffusivity with increased polysilicate loading. Perhaps most enlightening was the 
evidence that base catalysis produces polysilicate phases that are much more absorptive in 
nature than acid catalysis. Taken as a whole, the results point to sub-nanometer porosity 
within the base catalyzed inorganic domains and even smaller pores, essentially bonding 
defects, within the acid catalyzed polysilicate domains. 

An additional level of complexity was present in this study in that a luminescence- 
based technique was used to measure the diffusivity. While this technique is certainly not 
difficult, the influence of molecular environment upon the luminescence characteristics of 
the compound becomes important. This technique also limits the results somewhat in that 
no measure of permeability exists, and consequently, no method of calculating an average 
solubility exists. This is important because there is no way of checking the validity of the 
Henry's law assumption this luminescent technique relies upon. Traditional concentration- 
pressure studies utilizing a micro-balance or other technique need to be performed The 
insights gained from such sorption studies would undoubtedly provide a valuable second 
opinion of the absorptive capabilities of the inorganic phases in these gels. 

Another approach that warrants utilization is that of preferentially localizing the 
lumiphores within each phase. For example, many different types of luminescent 



153 

compounds exist. If a salt-based lumiphore were to be present during the prehydrolysis of 
TEOS, then it is likely that polar interactions between the silanol groups of the gelling 
polysilicate and lumiphore would result in incorporation of the probe into the glassy 
network. These doped polysilicate structures could then be processed with the end 
functionalized PTMO and gels cast. For the sake of comparison, the same probe could be 
dispersed into the PTMO sol and the prehyrolyzed TEOS added immediately prior to 
casting. This would force the lumiphore into the PTMO rich phase. By measuring the 
diffusivity of both types of gels as a function of polysilicate loading, the effects of 
environment upon diffusivity could be discerned and a check made of the premise that the 
lumiphore must be present in every phase for which it is assessing the diffusivity. 



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BIOGRAPHICAL SKETCH 

The author was born in St. Louis, Missouri, in 1969. His father, Wayne Miller, 
was a pilot in the United States Army. His father's occupation resulted in twenty-one 
changes of address throughout the eastern U.S. before the author was 10 years old. The 
families twenty-second move was to Paducah, Kentucky, where the author and parents 
began life as civilians. While growing up in Kentucky, the author enjoyed hobbies of 
fishing, target shooting and especially windsurfing on nearby Kentucky Lake. In 1987 he 
began his studies in ceramic engineering at the University of Missouri at Rolla. He 
graduated with a Bachelor of Science degree and became engaged to Yvonne Gander in 
the fall of 1991. He immediately began pursuing a master's degree in materials science at 
the University of Florida. In May of 1992 he married. His studies and research focused 
on the sol-gel synthesis and characterization of organic-inorganic hybrid composites. He 
earned the master's degree in December of 1993. With the exception of a two year long 
quest in search of the ultimate slow-to-gel, thermal shock resistant epoxy, his efforts 
continued in the field of hybrid technology. Upon receiving his doctorate, he will 
immediately begin working in industry, which he considers the reward for the many years 
invested in formal education. 









163 



I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 



£/%~ 



Anthony B. B/ennan, Chairman 
Associate Professor of Materials Science 
and Engineering 



I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 

Christopher D. Batich 

Professor of Materials Science and 

Engineering 

I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 




Elliot P. 

AssistanfProfessor of Materials Science 
and Engineering 

I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 





Eugene P! Goldberg 
Professor of Materials Science and N 
Engineering 









I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 



KjjihtjL 



8 U. 



Kenneth B. Wagener 
Professor of Chemistry 

This dissertation was submitted to the Graduate Faculty of the College of 
Engineering and to the Graduate School and was accepted as partial fulfillment of the 
requirements for the degree of Doctor of Philosophy. 



December, 1997 



£ 



Winfred M. Phillips 

Dean, College of Engineering 



Karen A. Holbrook 
Dean, Graduate School 






LD 

1780 

199.3 

M(p5I 



UNIVERSITY OF FLORIDA 



3 1262 08555 0902